High-ductility steel sheet excellent in press formability and strain age hardenability, and method for manufacturing the same

ABSTRACT

A steel sheet composition contains appropriate amounts of C, Si, Mn, P, S, Al and N and 0.5 to 3.0% Cu. A composite structure of the steel sheet has a ferrite phase or a ferrite phase and a tempered martensite phase as a primary phase, and a secondary phase containing retained austenite in a volume ratio of not less than 1%. In place of the Cu, at least one of Mo, Cr, and W may be contained in a total amount of not more than 2.0%. This composition is useful in production of a high-ductility hot-rolled steel sheet, a high-ductility cold-rolled steel sheet and a high-ductility hot-dip galvanized steel sheet having excellent press formability and excellent stain age hardenability as represented by a ΔTS of not less than 80 MPa, in which the tensile strength increases remarkably through a heat treatment at a relatively low temperature after press forming.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates mainly to steel sheets for automobiles,and more particularly, to high-ductility steel sheets having very highstrain age hardenability and excellent press formability such asductility, stretch-flanging formability, and drawability, in which thetensile strength increases remarkably through a heat treatment afterpress forming, and to methods for manufacturing the same. The term“steel sheets” as herein used shall include hot-rolled steel sheets,cold-rolled steel sheets, and hot-dip galvanized steel sheets. The term“steel sheets” as herein used shall also include steel sheets and steelstrips.

2. Description of the Related Art

In recent years, weight reduction in automobile bodies has become a veryimportant issue in relation to emission gas control for the purpose ofpreserving global environments. More recently, efforts are made toachieve higher strength of automotive steel sheets and to reduce steelsheet thickness in order to reduce the weights of automobile bodies.

Because most of the body parts of automobiles made of steel sheets areformed by press working, steel sheets used must have excellent pressformability. In order to achieve excellent press formability, it isnecessary to ensure high ductility. Stretch flanging is frequentlyapplied, so that the steel sheets to be used must have a highhole-expanding ratio. In general, however, a higher strength of steelsheet tends to result in a lower ductility and a lower hole-expandingratio, thus leading to poor press formability. As a result, there hasconventionally been an increasing demand for high-strength steel sheetshaving high ductility and excellent press formability.

Importance is now placed on safety of an automobile body to protect adriver and passengers upon collision, and for this purpose, steel sheetsmust have improved impact resistance as a standard of safety uponcollision. For the purpose of improving the crashworthiness, a higherstrength in a completed automobile is more favorable. There hastherefore been the strongest demand for steel sheets having lowstrength, high ductility, and excellent press formability upon formingautomobile parts, and having high strength and excellent crashworthinessin completed products.

To satisfy such a demand, a steel sheet high both in press formabilityand strength was developed. This is a bake hardenable type steel sheetof which the yield stress increases by applying a bake treatmentincluding holding at a high temperature of 100 to 200° C. after pressforming. In this steel sheet, the C content remaining finally in a solidsolution state (solute C content) is controlled within an appropriaterange so as to keep the softness, shape fixability, and ductility duringpress forming. In a bake treatment performed after the press forming ofthis steel sheet, the solute C is fixed to a dislocation introducedduring the press forming and inhibits the movement of the dislocation,thus resulting in an increase in yield stress. In this bake hardenabletype automotive steel sheet, the yield stress can be increased, but thetensile strength cannot be increased.

Japanese Examined Patent Application Publication No. 5-24979 discloses abake hardenable high-strength cold-rolled steel sheet having acomposition comprising C: 0.08 to 0.20%, Mn: 1.5 to 3.5% and the balanceFe and incidental impurities, and having a structure composed of uniformbainite containing not more than 5% of ferrite or composed of bainitepartially containing martensite. The cold-rolled steel sheet disclosedin Japanese Examined Patent Publication No. 5-24979 is manufactured byrapidly cooling the steel sheet to a temperature in the range of 400 to200° C. in the cooling step after continuous annealing and then slowlycooling the same. A high degree of baking hardening conventionallyunavailable is thereby achieved through conversion from the conventionalstructure mainly comprising ferrite to a structure mainly comprisingbainite in the steel sheet.

In the steel sheet disclosed in Japanese Examined Patent ApplicationPublication No. 5-24979, a high degree of baking hardeningconventionally unavailable is obtained through an increase in yieldstrength after bake treatment. Even in this steel sheet, however, it isyet difficult to increase tensile strength after the bake treatment, andan improvement in crashworthiness cannot still be achieved.

On the other hand, some hot-rolled steel sheets are proposed with a viewto increasing not only yield stress but also tensile strength byapplying a heat treatment after press forming.

For example, Japanese Examined Patent Application Publication No.8-23048 proposes a method for manufacturing a hot-rolled steel sheetcomprising the steps of reheating a steel containing C: 0.02 to 0.13%,Si: not more than 2.0%, Mn: 0.6 to 2.5%, sol. Al: not more than 0.10%,and N: 0.0080 to 0.0250% to a temperature of not less than 1,100° C. andapplying hot finish rolling at a temperature of 850 to 950° C. Themethod also comprising the steps of cooling the hot-rolled steel sheetat a cooling rate of not less than 15° C./second to a temperature ofless than 150° C., and coiling the same, thereby forming a compositestructure mainly comprising ferrite and martensite. In the steel sheetmanufactured by the technique disclosed in Japanese Examined PatentApplication Publication No. 8-23048, the tensile strength and the yieldstress increase by strain age hardening; however, a serious problem isposed in that coiling of the steel sheet at a very low coilingtemperature as less than 150° C. results in large variations inmechanical properties. Another problem includes a large variation inincrement of yield stress after press forming and bake treatments, aswell as poor press formability due to a low hole-expanding ratio (λ) anddecreased stretch-flanging workability.

Japanese Unexamined Patent Application Publication No. 11-199975proposes a hot-rolled steel sheet for working excellent in fatiguecharacteristics, containing C: 0.03 to 0.20%, appropriate amounts of Si,Mn, P, S and Al, Cu: 0.2 to 2.0%, and B: 0.0002 to 0.002%, of which themicrostructure is a composite structure comprising ferrite as a primaryphase and martensite as a second phase, and the ferrite phase containsCu in a solid-solution and/or precipitation state of not more than 2 nm.The steel sheet disclosed in Japanese Unexamined Patent ApplicationPublication No. 11-199975 has an object based on the fact that thefatigue limit ratio is remarkably improved only when Cu and B are addedin combination, and Cu is present in an ultra fine state not more than 2nm. For this purpose, it is essential to complete hot finish rolling ata temperature above the A_(r3) transformation point, air-cool the sheetwithin the temperature region of A_(r3) to A_(r1) for 1 to 10 seconds,cool the sheet at a cooling rate of not less than 20° C./second, andcoil the cooled sheet at a temperature of not more than 350° C. A lowcoiling temperature of not more than 350° C. causes serious deformationof the shape of the hot-rolled steel sheet, thus inhibiting industriallystable manufacture.

On the other hand, some automobile parts must have high corrosionresistance. A hot-dip galvanized steel sheet is suitable as a materialapplied to portions requiring high corrosion resistance. For thisreason, a particular demand exists for hot-dip galvanized steel sheetsexcellent in press formability during forming, and is considerablyhardened by a heat treatment after the forming.

To respond to such a demand, for example, Japanese Patent PublicationNo. 2802513 proposes a method for manufacturing a hot-dip galvanizedsteel sheet using a hot-rolled steel sheet as a black plate. The methodcomprises the steps of hot-rolling a steel slab containing C: not morethan 0.05%, Mn: 0.05 to 0.5%, Al: not more than 0.1% and Cu: 0.8 to 2.0%at a coiling temperature of not more than 530° C. The method furthercomprising the subsequent steps of reducing the steel sheet surface byheating the hot-rolled steel sheet to a temperature of not more than530° C., and hot-dip-galvanizing the sheet, whereby remarkable hardeningis available through a heat treatment after forming. In the steel sheetmanufactured by this method, however, the heat treatment temperaturemust be high as not less than 500° C., in order to obtain remarkablehardening from the heat treatment after the forming, and this has aproblem in practice.

Japanese Unexamined Patent Application Publication No. 10-310824proposes a method for manufacturing an alloyed hot-dip galvanized steelsheet having increased strength by a heat treatment after forming, usinga hot-rolled or cold-rolled steel sheet as a black plate. This methodcomprises the steps of hot-rolling a steel containing C: 0.01 to 0.08%,appropriate amounts of Si, Mn, P, S, Al and N, and at least one of Cr, Wand Mo: 0.05 to 3.0% in total. The method further comprises the step ofcold-rolling or temper-rolling and annealing the sheet. The method stillfurther comprises the step of applying hot-dip galvanizing to the sheetand heating the sheet for alloying treatment. The tensile strength ofthe steel sheet is increased by heating the sheet at a temperaturewithin the range of 200 to 450° C. However, the resultant steel sheetinvolves a problem in that the microstructure comprises a ferrite singlephase, a ferrite and pearlite composite structure, or a ferrite andbainite composite structure; hence, high ductility and low yieldstrength are unavailable, resulting in low press formability.

SUMMARY OF THE INVENTION

The present invention was made in view of the fact that, in spite of thestrong demand as described above, a technique for industrially stablymanufacturing a steel sheet satisfying these properties has never beenfound. The present invention solves the problems described above. It isan object of the present invention to provide is directed tohigh-ductility and high-strength steel sheets suitable for automobilesand having excellent press formability and excellent strain agehardenability, in which the tensile strength increases considerablythrough a heat treatment at a relatively low temperature after pressforming. It is also an object of the present invention to provide amanufacturing method capable of stably manufacturing the high-ductilityand high-strength steel sheets.

To achieve the above-mentioned object of the invention, the inventorscarried out extensive studies on the effect of the steel sheet structureand alloying elements on strain age hardenability. As a result, theinventors found that a steel sheet having high age hardenability whichleads to both an increase in yield stress and a remarkable increase intensile strength can be obtained after a pre-deformation treatment witha prestrain of not less than 5% and a heat treatment at a relatively lowtemperature as within the range of 150 to 350° C. by (1) forming acomposite structure of the steel sheet comprising ferrite and a phasecontaining retained austenite in a volume ratio of not less than 1%, and(2) limiting the C content within the range of a low-carbon region to amedium-carbon region and containing Cu within an appropriate range or atleast one of Mo, Cr, and W in place of Cu. In addition, the steel sheetwas found to have satisfactory ductility, a high hole expanding ratio,and excellent press formability.

The results of a fundamental experiment carried out by the inventors onhot-rolled steel sheets will first be described.

A sheet bar having a composition comprising, in weight percent, C:0.10%, Si: 1.4%, Mn: 1.5%, P: 0.01%, S: 0.005%, Al: 0.04%, N: 0.002% andCu: 0.3 or 1.3% was heated to 1,250° C. and soaked. Then, the sheet barwas subjected to three-pass rolling into a thickness of 2.0 mm so thatthe finish rolling end temperature was 850° C. Thereafter, coolingconditions and the coiling temperature were changed variously to converta single ferrite structure steel sheet into a hot-rolled steel sheetwith a composite structure composed of ferrite as a primary phase and aretained austenite-containing phase as a secondary phase (hereinafter,referred to also as a composite ferrite/retained austenite structure).

Tensile properties were investigated by a tensile test on the resultanthot-rolled steel sheets. A pre-deformation treatment of a tensileprestrain of 5% was applied to each test piece sampled from thesehot-rolled steel sheets. Then, after applying a heat treatment at 50 to350° C. for 20 minutes, a tensile test was carried out to determinetensile properties, and the strain age hardenability was evaluated.

The strain age hardenability was evaluated in terms of the increment ΔTSthat is a difference between the tensile strength TS_(HT) after heattreatment and the tensile strength TS before the heat treatment. Thatis, ΔTS=(tensile strength TS_(HT) after heat treatment)−(tensilestrength TS before pre-deformation treatment). The tensile test wascarried out by using JIS No. 5 tensile test pieces sampled in therolling direction.

FIG. 1 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the steel sheet structure. A pre-deformation treatmentof a tensile prestrain of 5% and then a heat treatment of 250° C. for 20minutes were applied to the test pieces. The increment ΔTS wasdetermined from the difference in tensile strength TS between before andafter the heat treatment. FIG. 1 suggests that, for a Cu content of 1.3wt. %, a high strain age hardenability as represented by a ΔTS of notless than 80 MPa is obtained by forming a composite ferrite/retainedaustenite steel sheet structure. For a Cu content of 0.3 wt. %, ΔTS isless than 80 MPa, irrespective of the steel sheet structure, and highstrain age hardenability cannot be obtained.

It is possible to manufacture a hot-rolled steel sheet having a highstrain age hardenability by limiting the Cu content within anappropriate range, and forming a composite structure having ferrite as aprimary phase and a retained austenite-containing phase as a secondaryphase.

FIG. 2 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the heat treatment temperature after pre-straintreatment. The microstructure of the steel sheet is a compositestructure having ferrite as a primary phase and a retainedaustenite-containing phase as a secondary phase, and the volume ratio ofthe retained austenite structure is 8% of the entire structure.

FIG. 2 shows that the increment ΔTS increases as the heat treatmenttemperature increases and strongly depends on the Cu content. With a Cucontent of 1.3 wt. %, a high strain age hardenability as represented bya ΔTS of not less than 80 MPa is obtained at a heat treatmenttemperature of not less than 150° C. For a Cu content of 0.3 wt. %, ΔTSis less than 80 MPa at any heat treatment temperature, and high strainage hardenability cannot be obtained.

In addition, a hole expanding test was carried out on steel sheetshaving a single ferrite structure or a composite ferrite/retainedaustenite structure, and Cu contents of 0.3 wt % and 1.3 wt %, and thehole expanding ratio λ was determined. In the hole expanding test, punchholes were formed in test pieces through punching with a punch having adiameter of 10 mm. Thereafter, hole expansion was conducted with aconical punch having a vertical angle of 60 degrees so that the burr wasoutside, until cracks passing through the sheet in the thicknessdirection form. The hole expanding ratio λ was determined by theformula: λ(%)={(d−d₀)/d₀}×100 where d₀ represents the initial holediameter, and d represents the hole inside diameter on occurrence ofcracks.

In the case of a Cu content of 1.3 wt %, a hot-rolled steel sheet havinga composite ferrite/retained austenite structure had a hole expandingratio of about 140%, and a hot-rolled steel sheet having a singleferrite structure also had a hole expanding ratio of about 140%. Incontrast, in the case of a Cu content of 0.3 wt %, a hot-rolled steelsheet having a single ferrite structure had a hole expanding ratio of120%, and a hot-rolled steel sheet having a composite ferrite/retainedaustenite structure had a hole expanding ratio of about 80%.

As described above, it is clear that the hot-rolled steel sheet having acomposite ferrite/retained austenite structure has an increased holeexpanding ratio and that hole expanding formability is improved with anincreased Cu content. A detailed mechanism of the improvement in holeexpanding formability by Cu has not yet been clarified. The contained Cuis considered to reduce the difference in hardness between theferrite/retained austenite and the strain-induced transformedmartensite.

In the hot-rolled steel sheet of the present invention, very fine Cuprecipitates in the steel sheet as a result of a pre-deformation with astrain of 2% or more as measured upon measuring the increment ofdeformation stress from before to after a usual heat treatment and theheat treatment carried out at a relatively low temperature in the rangeof 150 to 350° C. According to a study carried out by the presentinventors, high strain age hardenability bringing about an increase inyield stress and a remarkable increase in tensile strength probablyachieved by the precipitation of very fine Cu. Such precipitation ofvery fine Cu by a heat treatment in a low-temperature region has neverbeen observed in ultra-low carbon steel or low-carbon steel in reportsso far released. A reason for precipitation of very fine Cu in a heattreatment at a low temperature has not as yet been clarified to date.However, it is presumable as follows. During isothermal holding in thetemperature range of 620 to 780° C. or during slow cooling from thistemperature range after rapid cooling subsequent to hot rolling, a largeamount of Cu is distributed to the γ phase. After cooling, Cu isdissolved in the retained austenite in a supersaturation state. Theretained austenite is transformed into martensite by a prestrain of notless than 5%, and very fine Cu precipitates in the strain-inducedtransformed martensite during a subsequent low-temperature treatment.

Next, the results of a fundamental experiment carried out by the presentinventors on the cold-rolled steel sheet will be described.

A sheet bar having a composition comprising, in weight percent, C:0.10%, Si: 1.2%, Mn: 1.4%, P: 0.01%, S: 0.005%, Al: 0.03%, N: 0.002%,and Cu: 0.3 or 1.3% was heated to 1,250° C., soaked and subjected tothree-pass rolling into a thickness of 4.0 mm so that the finish rollingend temperature was 900° C. After the completion of finish rolling, atemperature holding equivalent treatment of 600° C. for 1 hour wasapplied as a coiling treatment. Thereafter, the sheet was cold-rolled ata reduction of 70% into a cold-rolled steel sheet having a thickness of1.2 mm. Then, the cold-rolled sheet was heated at a temperature in therange of 700 to 850° C. and soaked for 60 seconds. Thereafter, the sheetwas cooled to 400° C., and was held at the temperature (400° C.) for 300seconds for recrystallization annealing. By the recrystallizationannealing, various cold-rolled steel sheets were obtained in which thestructure changed from a single ferrite structure to a compositeferrite/retained austenite structure.

Tensile tests were conducted on the resultant cold-roll steel sheets asin the hot-rolled steel sheets to determine tensile properties. Tensileproperties (YS, TS) were determined by sampling test pieces from thesecold-rolled steel sheets, applying a pre-deformation treatment with atensile prestrain of 5% to these test pieces, then heating the steelsheets at 50 to 350° C. for 20 minutes, and then conducting the tensiletests.

The strain age hardenability was evaluated in terms of the tensilestrength increment ΔTS from before to after the heat treatment, as inthe hot-rolled steel sheet.

FIG. 3 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the recrystallization annealing temperature. The valueΔTS was determined by applying a pre-deformation treatment with atensile prestrain of 5% to test pieces sampled from the resultantcold-rolled steel sheets, conducting a heat treatment of 250° C. for 20minutes, and carrying out a tensile test.

FIG. 3 suggests that a high strain age hardenability as represented by aΔTS of not less than 80 MPa is available, in the case of a Cu content of1.3 wt. %, by employing a recrystallization annealing temperature of notless than 750° C. to convert the steel sheet structure into a compositeferrite/retained austenite structure. On the other hand, in the case ofa Cu content of 0.3 wt. %, high strain age hardenability is unavailablebecause ΔTS is less than 80 MPa at any recrystallization annealingtemperature. FIG. 3 suggests the possibility of manufacturing acold-rolled steel sheet having a high strain age hardenability byoptimizing the Cu content and forming a composite ferrite/retainedaustenite structure.

FIG. 4 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the heat treatment temperature after pre-straintreatment. The steel sheet used was annealed at 800° C., which was thedual phase region of ferrite (α)+austenite (γ), for a holding time of 60seconds after cold rolling, cooled from the holding temperature (800°C.) to 400° C. at a cooling rate of 30+ C./second, and held at 400° C.for 300 seconds. The steel sheets had a composite ferrite/retainedaustenite (secondary phase) microstructure, the volume ratio of theretained austenite structure being 4%.

FIG. 4 shows that the increment ΔTS increases as the heat treatmenttemperature increases and strongly depends on the Cu content. With a Cucontent of 1.3 wt. %, a high strain age hardenability as represented bya ΔTS of not less than 80 MPa is obtained at a heat treatmenttemperature of not less than 150° C. For a Cu content of 0.3 wt. %, ΔTSis less than 80 MPa at any heat treatment temperature, and high strainage hardenability cannot be obtained.

In addition, a hole expanding test was carried on cold-rolled steelsheets having a composite ferrite/retained austenite structure and Cucontents of 0.3 wt % and 1.3 wt. % to determine the hole expanding ratio(λ), as in the hot-rolled steel sheet.

In the cold-rolled steel sheet with a Cu content of 1.3%, λ was 130%;while in the cold-rolled steel sheet with a Cu content of 0.3%, λ was60%. It is clear that, for a Cu content of 1.3 wt. %, the hole expandingratio is increased and hole expanding formability is improved even inthe cold-rolled steel sheet, as in the hot-rolled steel sheet. Adetailed mechanism of improvement in hole expanding formability withcontent of Cu has not yet been clarified, as in the hot-rolled steelsheet. Also, in the cold-rolled steel sheet, it is considered that thecontained Cu reduces the difference in hardness between theferrite/retained austenite structure and the strain-induced transformedmartensite structure.

In the cold-rolled steel sheet of the present invention, very fine Cuprecipitates in the steel sheet as a result of a pre-deformation with astrain larger than 2%, which is equivalent to the prestrain on measuringthe deformation stress increment from before to after a usual heattreatment, and a heat treatment at a relatively low temperature of 150to 350° C. According to a study carried out by the present inventors,also in the cold-rolled steel sheet, high strain age hardenabilitybringing about an increase in yield stress and a remarkable increase intensile strength is probably achieved by the precipitation of very fineCu. A reason for precipitation of very fine Cu in a heat treatment in alow temperature region has not as yet been clarified to date. However,it is presumable as follows. During recrystallization annealing in thedual phase region of α+γ, a large amount of Cu is distributed to the γphase. The distributed Cu remains even after cooling and is dissolvedinto the martensite in a supersaturation state, and very fine Cuprecipitates through a prestrain of not less than 5% and alow-temperature treatment.

Next, the result of a fundamental experiment carried out by the presentinventors on the hot-dip galvanized steel sheet will be described.

A sheet bar having a composition comprising, in weight percent, C:0.08%, Si: 0.5%, Mn: 2.0%, P: 0.01%, S: 0.004%, Al: 0.04%, N: 0.002% andCu: 0.3 or 1.3% was heated to 1,250° C. and soaked. Then, the sheet barwas subjected to three-pass rolling into a thickness of 4.0 mm so thatthe finish rolling end temperature was 900° C. After the finish rolling,a temperature holding equivalent treatment of 600° C. for 1 h wasapplied as a coiling treatment. Thereafter, the hot-rolled sheet wascold-rolled at a reduction of 70% into a cold-rolled steel sheet havinga thickness of 1.2 mm. Then, the cold-rolled sheet was heated and soakedat 900° C., and cooled at a cooling rate of 30° C./sec. (a primary heattreatment). The steel sheet after the primary heat treatment had a lathmartensite structure. The steel sheet after the primary heat treatmentwas subjected to a secondary heat treatment at various temperatures,then rapidly cooled to a temperature in the range of 450 to 500° C.Then, the sheet was immersed into a hot-dip galvanizing bath (0.13 wt. %Al—Zn bath) to form a hot-dip galvanizing layer on the surface. Further,the sheet was reheated to a temperature in the range of 450 to 550° C.to alloy the hot-dip galvanizing layer (Fe content in the galvanizinglayer: about 10%).

For the resultant hot-dip galvanized steel sheet, tensile propertieswere determined through a tensile test. In addition, test pieces weresampled from the hot-dip galvanized steel sheet, and a pre-deformationtreatment with a tensile prestrain of 5% was applied to the test pieces,as in the hot-rolled steel sheet and the cold-rolled steel sheet. Then,a heat treatment of 50 to 350° C. for 20 minutes was applied.Thereafter, a tensile test was carried out to determine tensileproperties. The strain age hardenability was evaluated in terms of theincrement ΔTS of the tensile strength from before to after the heattreatment.

FIG. 5 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the secondary heat treatment temperature. The incrementΔTS was determined by applying a tensile prestrain of 5% to test piecessampled from the resultant hot-dip galvanized steel sheets, conducting aheat treatment at 250° C. for 20 minutes, and carrying out a tensiletest.

FIG. 5 suggests that, for a Cu content of 1.3 wt. %, a high strain agehardenability as represented by a ΔTS of not less than 80 MPa can beobtained by forming a composite ferrite/tempered martensite/retainedaustenite steel sheet structure. In contrast, in the case of a Cucontent of 0.3 wt. %, high strain age hardenability cannot be obtainedas because ΔTS is less than 80 MPa at any secondary heat treatmenttemperature.

FIG. 5 suggests the possibility of manufacturing a hot-dip galvanizedsteel sheet having high strain age hardenability by optimizing the Cucontent and by forming a composite ferrite/tempered martensite/retainedaustenite structure.

FIG. 6 illustrates the effect of the Cu content on the relationshipbetween ΔTS and the heat treatment temperature after pre-straintreatment. The increment ΔTS was determined by applying a tensileprestrain of 5% to test pieces sampled from the alloyed hot-dipgalvanized steel sheets treated at a secondary heat treatmenttemperature of 800° C., conducting a heat treatment of 50 to 350° C. for20 minutes, and carrying out a tensile test.

FIG. 6 shows that the increment ΔTS increases as the heat treatmenttemperature increases after the pre-deformation treatment and stronglydepends on the Cu content. With a Cu content of 1.3 wt. %, a high strainage hardenability as represented by a ΔTS of not less than 80 MPa can beobtained at a heat treatment temperature of not less than 150° C. Incontrast, for a Cu content of 0.3 wt. %, ΔTS is less than 80 MPa at anyheat treatment temperature, and high strain age hardenability cannot beobtained.

In the hot-dip galvanized steel sheet of the present invention, veryfine Cu precipitates in the steel sheet as a result of a pre-deformationwith a strain larger than 2% which is a usual amount of strain onmeasuring the deformation stress increment from before to after a heattreatment, and a heat treatment within a relatively low temperatureregion of 150 to 350° C. According to a study carried out by the presentinventors, high strain age hardenability bringing about an increase inyield stress and a remarkable increase in tensile strength is probablyachieved by the precipitation of very fine Cu. A reason forprecipitation of very fine Cu in a heat treatment in a low temperatureregion has not as yet been clarified to date. However, it is presumableas follows. During heat treatment in the dual phase region of ferrite(α)+austenite (γ), a large amount of Cu is distributed to the γ phase,and the distributed Cu remaining even after cooling is dissolved intothe retained austenite in a supersaturation state. The retainedaustenite is transformed into martensite by a prestrain of not less than5%, and very fine Cu precipitates in the martensite through a subsequentlow-temperature heat treatment.

In addition, hole expanding test was performed using hot-dip galvanizedsteel sheets having a composite structure of ferrite/temperedmartensite/retained austenite and Cu contents of 0.3 wt % and 1.3 wt. %to determine the hole expanding ratio (λ), as in the hot-rolled steelsheet and the cold-rolled steel sheet.

The hole expanding ratio λ of the steel sheet having a Cu content of1.3% was 120%, while the hole expanding ratio λ of the steel sheethaving a Cu content of 0.3% was 50%. The results suggest that for a Cucontent of 1.3 wt %, the hole expanding ratio is increased and holeexpanding formability is improved, as compared with a Cu content of0.3%.

A detailed mechanism of improvement in hole expanding formability withcontent of Cu has not yet been clarified, as in the hot-rolled steelsheet and the cold-rolled steel sheet, but it is considered that thecontained Cu reduces the difference in hardness among the ferrite, thetempered martensite/retained austenite, and the martensite formed bystrain induced transformation.

On the basis of the novel findings as described above, the presentinventors carried out further extensive studies and found that theabove-mentioned phenomena occurred in steel sheets not containing Cu aswell.

The structure of a steel sheet having a composition containing at leastone of Mo, Cr, and W was converted to a composite structure containing aferrite primary phase and a phase containing retained austenite as asecondary phase. Thereafter, by applying a prestrain and a heattreatment in a low temperature region, it was found that very finecarbides precipitated in the strain-induced transformed martensite,resulting in an increase in tensile strength. The strain-induced fineprecipitation at a low temperature was more remarkable in a steelcomposition containing at least one of Nb, Ti, and V in addition to atleast one of Mo, Cr, and W.

The present invention was completed through further studies on the basisof the aforementioned findings. The gist of the present invention is asfollows:

-   (1) A high-ductility steel sheet excellent in press formability and    in strain age hardenability as represented by a ΔTS of not less than    80 MPa, comprising a composite structure containing a primary phase    containing a ferrite phase and a secondary phase containing a    retained austenite phase in a volume ratio of not less than 1%.-   (2) A high-ductility steel sheet according to aspect (1), wherein    the steel sheet is a hot-rolled steel sheet, and the primary phase    consisting essentially of a ferrite phase.-   (3) A high-ductility steel sheet according to aspect (2), wherein    the hot-rolled steel sheet has a composition comprising, in weight    percent, C: 0.05 to 0.20%, Si: l.0 to 3.0%, Mn: not more than 3.0%,    P: not more than 0.10%, S: not more than 0.02%, Al: not more than    0.30%, N: not more than 0.02%, and Cu: 0.5 to 3.0%, and the balance    Fe and incidental impurities.-   (4) A high-ductility steel sheet according to aspect (3), the    composition further comprising, in weight percent, at least one of    the following Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (5) A high-ductility steel sheet according to aspect (2), wherein    the hot-rolled steel sheet has a composition comprising, in weight    percent, C: 0.05 to 0.20%, Si: 1.0 to 3.0%, Mn: not more than 3.0%,    P: not more than 0.10%, S: not more than 0.02%, Al: not more than    0.30%, N: not more than 0.02%, at least one of Mo: 0.05 to 2.0%, Cr:    0.05 to 2.0% and W: 0.05 to 2.0%, not more than 2.0% in total, and    the balance Fe and incidental impurities.-   (6) A high-ductility steel sheet according to aspect (5), the    composition further containing, in weight percent, at least one of    Nb, Ti, and V in an amount of not more than 2.0% in total.-   (7) A method for manufacturing a high-ductility hot-rolled steel    sheet excellent in press formability and in strain age hardenability    as represented by a ΔTS of not less than 80 MPa, comprising the    steps of: hot-rolling a steel slab having a composition comprising,    in weight percent, C: not more than 0.20%, Si: 1.0 to 3.0%, Mn: not    more than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al:    not more than 0.30%, N: not more than 0.02%, and Cu: 0.5 to 3.0%,    into a hot-rolled steel sheet having a prescribed thickness, the hot    rolling step including finish-rolling at a finish rolling end    temperature of 780 to 980° C.; cooling the finish-rolled steel sheet    to a temperature in the range of 620 to 780° C. within 2 seconds at    a cooling rate of at least 50° C/second; holding the sheet at the    temperature in the range of 620 to 780° C. for 1 to 10 seconds, or    slowly cooling the sheet at a cooling rate of not more than 20°    C./second; cooling the sheet at a cooling rate of not less than 50°    C./second to a temperature of 300 to 500° C.; and coiling the sheet.-   (8) A method for manufacturing a high-ductility hot-rolled steel    sheet excellent in press formability and in strain age hardenability    as represented by a ΔTS of at least 80 MPa, according to aspect (7),    the composition further comprising, in weight percent, at least one    of the following Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (9) A method for manufacturing a high-ductility hot-rolled steel    sheet according to aspect (7), wherein the steel slab is replaced    with a steel slab having a composition containing, in weight    percent, C: 0.05 to 0.20%, Si: 1.0 to 3.0%, Mn: not more than 3.0%,    P: not more than 0.10%, S: not more than 0.02%, Al: not more than    0.30%, N: not more than 0.02%, and at least one of Mo: 0.05 to 2.0%,    Cr: 0.05 to 2.0% and W: 0.05 to 2.0% in a total amount of not more    than 2.0%.-   (10) A method for manufacturing a high-ductility hot-rolled steel    sheet according to aspect (9), the composition further containing,    in weight percent, at least one of Nb, Ti, and V in a total amount    of not more than 2.0%.-   (11) A method for manufacturing a high-ductility hot-rolled steel    sheet according to any one of aspects (7) to (10), wherein all or    part of the finish rolling is lubrication rolling.-   (12) A high-ductility steel sheet according to aspect (1), wherein    the steel sheet is a cold-rolled steel sheet, and the primary phase    containing the ferrite phase is a ferrite phase.-   (13) A high-ductility steel sheet according to aspect (12), wherein    the cold-rolled steel sheet has a composition comprising, in weight    percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not    more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:    not more than 0.3%, N: not more than 0.02%, Cu: 0.5 to 3.0%, and the    balance Fe and incidental impurities.-   (14) A high-ductility steel sheet according to aspect (13), the    composition further comprising, in weight percent, at least one of    the following Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (15) A high-ductility steel sheet according to aspect (12), wherein    the cold-rolled steel sheet has a composition comprising, in weight    percent: C: not more than 0.20%, Si: not more than 2.0%, Mn: not    more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:    not more than 0.3%, N: not more than 0.02%, at least one selected    from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0% and    W: 0.05 to 2.0%, not more than 2.0% in total, and the balance Fe and    incidental impurities.-   (16) A high-ductility steel sheet according to aspect (15), the    composition further comprising, in weight percent, at least one of    Nb, Ti, and V, in a total amount of not more than 2.0%.-   (17) A method for manufacturing a high-ductility cold-rolled steel    sheet excellent in press formability and in strain age hardenability    as represented by a ΔTS of not less than 80 MPa, comprising: a hot    rolling step of hot-rolling a steel slab having a composition    containing, in weight percent, C: not more than 0.20%, Si: not more    than 2.0%, Mn: not more than 3.0%, P: not more than 0.1%, S: not    more than 0.02%, Al: not more than 0.3%, N: not more than 0.02%, and    Cu: 0.5 to 3.0% as a material to form a hot-rolled steel sheet; a    cold rolling step of cold-rolling the hot-rolled steel sheet into a    cold-rolled steel sheet; and a recrystallization annealing step of    applying recrystallization annealing to the cold-rolled steel sheet    into a cold-rolled annealed steel sheet, the recrystallization    annealing step including a heat treatment of heating and soaking the    steel sheet in a ferrite/austenite dual phase region within a    temperature range of the A_(C1) transformation point to the A_(C3)    transformation point., cooling the sheet, and holding the sheet in    the temperature region of 300 to 500° C. for 30 to 1,200 seconds.-   (18) A method for manufacturing a high-ductility cold-rolled steel    sheet according to aspect (17), the composition further containing,    in weight percent, at least one selected from the following Groups A    to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (19) A method for manufacturing a high-ductility cold-rolled steel    sheet according to aspect (17), wherein the steel slab is replaced    with a steel slab having a composition containing, in weight    percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not    more than 3.0%, P: not more than 0.10%, S: not more than 0.02%, Al:    not more than 0.3%, N: not more than 0.02%, and at least one    selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to    2.0% and W: 0.05 to 2.0% in a total amount of not more than 2.0%.-   (20) A method of manufacturing a high-ductility cold-rolled steel    sheet according to aspect (19), the composition further containing,    in weight percent, at least one of Nb, Ti, and V in a total amount    of not more than 2.0%.-   (21) A method for manufacturing a high-ductility cold-rolled steel    sheet according to any one of aspects (17) to (20), wherein the    hot-rolling step includes heating the steel slab at a temperature of    not less than 900° C., rolling the slab at a finish rolling end    temperature of not less than 700° C., and coiling the hot-rolled    steel sheet at a coiling temperature of not more than 800° C.-   (22) A method for manufacturing a cold-rolled steel sheet according    to any one of aspects (17) to (21), wherein all or part of the hot    rolling is lubrication rolling.-   (23) A high-ductility hot-dip galvanized steel sheet comprising a    hot-dip galvanizing layer or an alloyed hot-dip galvanizing layer    formed on the surface of the high-ductility steel sheet according to    any one of aspects (1) to (6).-   (24) A high-ductility hot-dip galvanized steel sheet comprising a    hot-dip galvanizing layer or an alloyed hot-dip galvanizing layer    formed on the surface of the high-ductility steel sheet according to    any one of aspects (12) to (16).-   (25) A high-ductility steel sheet according to aspect (1), wherein    the steel sheet is a hot-dip galvanized steel sheet having a hot-dip    galvanizing layer or an alloyed hot-dip galvanizing layer formed on    a surface of the steel sheet, and the primary phase containing a    ferrite phase comprises a ferrite phase and a tempered martensite    phase.-   (26) A high-ductility steel sheet according to aspect (25), wherein    the steel sheet has a composition comprising, in weight percent, C:    not more than 0.20%, Si: not more than 2.0%, Mn: not more than 3.0%,    P: not more than 0.1%, S: not more than 0.02%, Al: not more than    0.3%, N: not more than 0.02%, Cu: 0.5 to 3.0%, and the balance Fe    and incidental impurities.-   (27) A high-ductility steel sheet according to aspect (26), the    composition further containing, in weight percent, at least one of    the following Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (28) A high-ductility steel sheet according to aspect (25), wherein    the steel sheet has a composition comprising, in weight percent, C:    not more than 0.20%, Si: not more than 2.0%, Mn: not more than 3.0%,    P: not more than 0.1%, S: not more than 0.02%, Al: not more than    0.3%, N: not more than 0.02%, at least one selected from the group    consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to 2.0%    in a total amount of not more than 2.0%, and the balance Fe and    incidental impurities.-   (29) A high-ductility steel sheet according to aspect (28), the    composition further containing, in weight percent, at least one of    Nb, Ti, and V in a total amount of not more than 2.0%.-   (30) A method for manufacturing of a high-ductility hot-dip    galvanized steel sheet excellent in press formability and in strain    age hardenability as represented by a ΔTS of not less than 80 MPa,    comprising: a primary heat-treating step of heating a steel sheet to    a temperature of not less than the A_(C1) transformation point and    rapidly cooling the steel sheet, the steel sheet having a    composition containing, in weight percent, C: not more than 0.20%,    Si: not more than 2.0%, Mn: not more than 3.0%, P: not more than    0.1%, S: not more than 0.02%, Al: not more than 0.3%, N: not more    than 0.02%, and Cu: 0.5 to 3.0%; a secondary heat-treating step of    heating the steel sheet to a temperature in the range of the A_(C1)    transformation point to the A_(C3) transformation point; and a    hot-dip galvanizing step of forming a hot-dip galvanizing layer on    the surface of the steel sheet.-   (31) A method for manufacturing a high-ductility cold-rolled steel    sheet according to aspect (30), the composition further containing,    in weight percent, at least one of the following Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

-   (32) A method for manufacturing a high-ductility hot-dip galvanized    steel according to aspect (30), wherein the steel sheet is replaced    with a steel sheet having a composition comprising, in weight    percent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not    more than 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al:    not more than 0.3%, N: not more than 0.02%, and at least one    selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to    2.0% and W: 0.05 to 2.0% in a total amount of not more than 2.0%.-   (33) A method for manufacturing a high-ductility hot-dip galvanized    steel sheet according to aspect (32), the composition further    containing, in weight percent, at least one of Nb, Ti, and V in a    total amount of not more than 2.0%.-   (34) A method for manufacturing a high-ductility hot-dip galvanized    steel sheet according to any one of aspects (30) to (33), further    comprising a pickling treatment step of pickling the steel sheet    between the primary heat-treating step and the secondary    heat-treating step.-   (35) A method for manufacturing a high-ductility hot-dip galvanized    steel sheet according to any one of aspects (30) to (34), further    comprising an alloying step of alloying the hot-dip galvanizing    layer, subsequent to the hot-dip galvanizing step.-   (36) A method for manufacturing a high-strength hot-dip galvanized    steel sheet according to any one of aspects (30) to (35), wherein    the steel sheet is a hot rolled steel sheet manufactured by    hot-rolling a material under conditions including a heating    temperature of not less than 900° C., a finish rolling end    temperature of not less than 700° C. and a coiling temperature of    not more than 800° C., or a cold-rolled steel sheet obtained by    cold-rolling the hot-rolled steel sheet.-   (37) A method for manufacturing a high-strength hot-dip galvanized    steel sheet according to aspect (36), wherein the cold-rolling is    performed at a reduction ratio of not less than 40%.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the steel sheet structure after apre-deformation and a heat treatment of a hot-rolled steel sheet;

FIG. 2 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the heat treatment temperature after apre-deformation and a heat treatment of a hot-rolled steel sheet;

FIG. 3 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the recrystallization annealing temperatureafter pre-deformation and a heat treatment of a cold-rolled steel sheet;

FIG. 4 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the heat treatment temperature afterpre-deformation and a heat treatment of a cold-rolled steel sheet;

FIG. 5 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the secondary heat treatment temperatureafter a pre-deformation and a heat treatment of a hot-dip galvanizedsteel sheet; and

FIG. 6 is a graph illustrating the effect of the Cu content on therelationship between ΔTS and the heat treatment temperature after apre-deformation and a heat treatment of a hot-dip galvanized steelsheet.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

A high-ductility steel sheet of the present invention has a tensilestrength TS of not less than 440 MPa, a composite structure comprising aprimary phase containing a ferrite phase and a secondary phasecontaining a retained austenite phase with a volume ratio of not lessthan 1%, excellent press formability, and excellent strain agehardenability, which is indicated by a remarkably increased tensilestrength ΔTS of not less than 80 MPa during a heat treatment at arelatively low temperature after press forming. The term “primary phase”used in the present invention shall be a structure occupying not lessthan 50% by a volume ratio.

The term “high-ductility steel sheet” used in the present inventionshall mean that a steel sheet has a balance (TS×El) of a tensilestrength (TS) and an elongation (El) of not less than 19,000 MPa %.

In addition, the term “ΔTS” used in the present invention means anincrement in tensile strength between before and after the heattreatment at a temperature in the range of 150 to 350° C. for a holdingtime of not less than 30 seconds of a steel sheet which was subjected toa pre-deformation treatment of a tensile plastic strain of not less than5%. That is, ΔTS=(tensile strength after heat treatment)−(tensilestrength before pre-deformation treatment). The steel sheets of thepresent invention shall include hot-rolled steel sheets, cold-rolledsteel sheets and hot-dip galvanized steel sheets.

All the steel sheets (hot-rolled steel sheets, cold-rolled steel sheetsand hot-dip galvanized steel sheets) having the above-mentionedstructure have high-ductility, excellent press formability, andexcellent strain age hardenability.

The term “superior strain age hardenability” or the term “excellentstrain age hardenability” used in the present invention shall mean that,when a steel sheet is subjected to a pre-deformation treatment of atensile plastic strain of not less than 5%, and then, to a heattreatment at a temperature in the range of 150 to 350° C. for a holdingtime of not less than 30 seconds, the increment ΔTS in tensile strengthbetween before and after the heat treatment is not less than 80 MPa,wherein ΔTS=(tensile strength TS_(HT) after heat treatment)−(tensilestrength TS before pre-deformation treatment). Preferably, the incrementΔTS is not less than 100 MPa. The heat treatment causes an increase ΔYSin yield stress of not less than 80 MPa, wherein ΔYS=(yield stressYS_(HT) after heat treatment)−(yield stress YS before pre-deformationtreatment).

In the control of the strain age hardenability, the amount of prestrain(pre-deformation) plays an important role. The present inventorsinvestigated the effect of the amount of prestrain on the subsequentstrain age hardenability by assuming possible deformation types appliedto automotive steel sheets. The results show that the uniaxialequivalent strain (tensile strain) is generally useful for elucidatingthe deformation of the steel sheets except for very deep drawing, thatthe uniaxial equivalent strain is mostly more than 5% for actual parts,and that the strength of the parts exhibit good correspondence to thestrength obtained after a strain aging treatment of a prestrain of 5%.Based on these findings, a tensile plastic strain of not less than 5% isemployed in the present invention.

Conventional bake treatment conditions include 170° C.×20 minutes as astandard. If precipitation strengthening by very fine Cu or fine carbideis performed as in the present invention, the heat treatment temperaturemust be 150° C. or more. Under conditions including a temperatureexceeding 350° C., on the other hand, the strengthening effect issaturated, and the steel sheet tends to soften. Heating to a temperatureexceeding 350° C. causes marked occurrence of thermal strain or tempercolor. For these reasons, a heat treatment temperature in the range of150 to 350° C. is adopted for strain age hardening in the presentinvention. The holding time of the heat treatment temperature should beat least 30 seconds. Holding a heat treatment temperature in the rangeof 150 to 350° C. for about 3.0 seconds permits achievement ofsubstantially satisfactory strain age hardening. For further enhancedstrain age hardening, the holding time is preferably at least 60seconds, and more preferably at least 300 seconds.

The heat treatment method after the pre-deformation is not limited inthe present invention, and atmospheric heating in a furnace in generalbake treatment, induction heating, non-oxidizing flame heating, laserheating, and plasma heating are suitably applicable. So-called hotpressing for pressing a heated steel sheet is also very effective meansin the present invention.

Next, the hot-rolled steel sheet, the cold-rolled steel sheet, and thehot-dip galvanized steel sheet in the present invention will bedescribed individually.

(1) Hot-rolled Steel Sheet

The hot-rolled steel sheet of the present invention will now bedescribed.

The hot-rolled steel sheet of the present invention has a compositestructure comprising a ferrite primary phase and a secondary phasecontaining a retained austenite phase having a volume ratio of not lessthan 1% of the entire structure. As described above, a hot-rolled steelsheet having such a composite structure exhibits high ductility, highstrength-ductility balance (TS×El), and excellent press formability.

Ferrite primary phase is preferably present in a volume ratio of notless than 50%. With a ferrite phase of less than 50%, it is difficult tokeep high ductility, resulting in lower press formability. When furtherenhanced ductility is required, the volume ratio of the ferrite phase ispreferably not less than 80%. For the purpose of making full use ofadvantages of the composite structure, the ferrite phase is preferablynot more than 98%.

In the present invention, steel must contain retained austenite phase asthe secondary phase in a volume ratio of not less than 1% of the entirestructure. With a retained austenite phase of less than 1%, highelongation (El) cannot be obtained. To obtain higher elongation (El),the retained austenite phase content is preferably not less than 2% andmore preferably not less than 3%.

The secondary phase may be a single retained austenite phase having avolume ratio of not less than 1%, or may be a mixture of a retainedaustenite phase of a volume ratio of not less than 1% and another phase,i.e., a pearlite phase, a bainite phase, and/or a martensite phase.

The reasons for limiting the composition of the hot-rolled steel sheetof the present invention will now, be described. The weight percent inthe composition will hereafter be denoted simply as %.

C: 0.05 to 0.20%

C is an element, which improves strength of a steel sheet and promotesthe formation of a composite structure of ferrite and retainedaustenite, and is preferably contained in an amount of not less than0.05% for forming the composite structure according to the presentinvention. A C content exceeding 0.20% causes an increase in proportionsof carbides in steel, resulting in a decrease in ductility, and hence adecrease in press formability. A more serious problem is that a Ccontent exceeding 0.20% leads to significant deterioration of spotweldability and arc weldability. For these reasons, the C content islimited within the range of 0.05 to 0.20% in the present invention. Fromthe viewpoint of formability, the C content is preferably not more than0.18%.

Si: 1.0 to 3.0%

Si is a useful strengthening element, which improves the strength of asteel sheet without a marked decrease in ductility of the steel sheet.In addition, Si is necessary for forming a retained austenite phase. Toobtain these effects, Si is preferably contained in an amount of notless than 1.0% and more preferably not less than 1.2%. An Si contentexceeding 3.0% leads to deterioration of press formability and degradesthe surface quality. The Si content is therefore limited within therange of 1.0 to 3.0%.

Mn: not more than 3.0%

Mn is a useful element, which strengthens steel and prevents hotcracking caused by S, and is therefore contained in an amount accordingto the S content. These effects are particularly remarkable at an Mncontent of not less than 0.5%. On the other hand, an Mn contentexceeding 3.0% results in deterioration of press formability andweldability. The Mn content is therefore limited to not more than 3.0%in the present invention. More preferably, the Mn content is not lessthan 1.0%.

P: not more than 0.10%

P strengthens steel, and may be contained in an amount necessary for adesired strength. From the viewpoint of increasing the strength, P ispreferably contained in an amount of not less than 0.005%. On the otherhand, a P content exceeding 0.10% results in deterioration of pressformability. The P content is therefore limited to not more than 0.10%.When superior press formability is required, the P content is preferablynot more than 0.08%.

S: not more than 0.02%

S is an element, which is present as inclusions in a steel sheet andcauses deterioration of ductility, formability, and particularly stretchflanging formability of the steel sheet, and it should be the lowestpossible. A reduced S content of not more than 0.02% does not exert muchadverse effect and therefore, the S content is limited to up to 0.02% inthe present invention. When more excellent stretch flanging formabilityis required, the S content is preferably not more than 0.010%.

Al: not more than 0.30%

Al is a useful element, which is added as a deoxidizing element tosteel, and improves cleanliness of steel. In addition, Al facilitatesthe formation of the retained austenite. These effects are particularlyremarkable at an Al content of not less than 0.01%. The Al contentexceeding 0.30% cannot give further effects, but causes deterioration ofpress formability. The Al content is therefore limited to not more than0.30%. Preferably, the Al content is not more than 0.10%. The presentinvention does not exclude a steelmaking process based on deoxidationusing a deoxidizer other than Al. For example, Ti deoxidation or Sideoxidation may be employed, and steel sheets produced by suchdeoxidation methods are also included in the scope of the presentinvention. In this case, addition of Ca or REM to molten steel does notimpair the features of the steel sheet of the present invention at all.

N: not less than 0.02%

N is an element, which increases the strength of a steel sheet throughsolid solution strengthening or strain age hardening, and is preferablycontained in an amount of not less than 0.0010% to obtain these effects.However, an N content exceeding 0.02% causes an increase in the contentof nitrides in the steel sheet, which causes serious deterioration ofductility, and thus, of press formability of the steel sheet. The Ncontent is therefore limited to not more than 0.02%. When furtherimprovement in press formability is required, the N content ispreferably not more than 0.01%, and more preferably less than 0.0050%.

Cu: 0.5 to 3.0%

Cu is an element, which remarkably increases strain age hardening of asteel sheet (increase in strength after pre-deformation/heat treatment),and thus is most important in the present invention. With a Cu contentof less than 0.5%, an increment ΔTS in tensile strength exceeding 80 MPacannot be obtained by changing the pre-determination/heat treatmentconditions. With a Cu content exceeding 3.0%, the effect is saturated sothat an effect corresponding to the content cannot be expected, leadingto unfavorable economic effects. Furthermore, deterioration of pressformability occurs, and the surface quality of the steel sheet isdegraded. The Cu content is therefore limited within a range of 0.5 to3.0%. In order to simultaneously achieve a higher ΔTS and excellentpress formability, the Cu content is preferably within a range of 1.0 to2.5%.

The hot-rolled steel sheet of the present invention containing Cupreferably further contains, in weight percent, at least one of thefollowing Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

Group A: Ni: not more than 2.0% Group A: Ni is effective for preventingthe formation of surface defects on the steel sheet surface containingCu, and may be added as required. The Ni content is preferably about ahalf the Cu content, i.e., in the range of about 30 to about 80% of theCu content. An Ni content exceeding 2.0% cannot give further enhancementin the effect because saturation of the effect, leading to economicdisadvantages, and causes deterioration of press formability. For thesereasons, the Ni content is preferably limited to not more than 2.0%.

Group B: at least one of Cr and Mo: not more than 2.0% in total

Group B: Both Cr and Mo, as well as Mn, strengthen the steel sheet andat least one thereof can be contained as required. This effect isparticularly remarkable at a Cr content of not less than 0.1% and at anMo content of not less than 0.1%. It is therefore preferable to containat least one of Cr: not less than 0.1% and Mo: not less than 0.1%. If atleast one of Cr and Mo are contained in a total amount exceeding 2.0%,press formability is impaired. It is therefore preferable to limit thetotal content of Cr and Mo to not more than 2.0%.

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total

Group C: Nb, Ti, and V are carbide-forming elements and effectivelyincrease the strength by fine dispersion of carbides, and can beselected and contained as required. This effect can be achieved at an Nbcontent of not less than 0.01%, a Ti content of not less than 0.01%, anda V content of not less than 0.01%. However, a total content of Nb, Ti,and V exceeding 0.2% causes deterioration of press formability. Thus,the total content of Nb, Ti, and V is preferably limited to not morethan 0.2%.

In the present invention, in place of the aforementioned Cu or at leastone of the above-mentioned Groups A to C, at least one selected from thegroup consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0%, and W: 0.05 to2.0% may be contained in an amount of not more than 2.0% in total, andat least one selected from the group consisting of Nb, Ti, and V may befurther contained in an amount of not more than 2.0% in total.

At least one selected from the group consisting of Mo: 0.05 to 2.0%, Cr:0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not more than 2.0% intotal

Mo, Cr, and W are elements, which remarkably increase strain agehardening (increase in strength after pre-deformation and heattreatment) of a steel sheet, and are one of the most important elementsin the present invention. That is, in the present invention, ahot-rolled steel sheet having a composite structure containing ferriteas a primary phase and a secondary phase of retained austenite andcontaining at least one of Mo, Cr, and W, causes strain-inducedtransformation of the retained austenite into martensite when aprestrain of not less than 5% and a low-temperature heat treatment areapplied to the hot-rolled steel sheet, and strain-induced fineprecipitation of fine carbides at a low temperature occurs in thestrain-induced transformed martensite, resulting in an increase intensile strength ΔTS of not less than 80 MPa. With a content of at leastone of Mo, Cr, and W of less than 0.05%, changing the steel sheetstructure and pre-deformation and heat treatment conditions does notcause an increase in tensile strength ΔTS of not less than 80 MPa. Onthe other hand, a content of at least one of Mo, Cr, and W exceeding2.0% does not give a corresponding effect because of saturation of theeffect, leading to economic disadvantages, and causes deterioration ofpress formability. The contents of Mo, Cr, and W are each preferablylimited within the range of 0.05 to 2.0%. From the viewpoint of pressformability, the total content of Mo, Cr and/or W is more preferablylimited to not more than 2.0%.

At least one of Nb, Ti, and V, in a total amount of not more than 2.0%

Nb, Ti, and V are carbide-forming elements, and can be added asrequired. Containing at least one of Nb, Ti, and V, in addition to atleast one of Mo, Cr, and W, and forming a composite structure containinga ferrite primary phase and a secondary phase of retained austenite formfine carbides in the strain-induced transformed martensite and causestrain-induced precipitation at low temperature, resulting in anincrease in tensile strength ΔTS of not less than 80 MPa. In order toobtain these effects, an Nb content is preferably not less than 0.01%, aTi content is preferably not less than 0.01%, and a V content ispreferably not less than 0.01%, and at least one of Nb, Ti, and V can beadded as required. However, a total content exceeding 2.0% causesdeterioration of press formability. Thus, the total content of Nb, Ti,and V is preferably limited to not more than 2.0%.

Apart from the above-mentioned elements, at least one of Ca: not lessthan 0.1% and REM: not less than 0.1% may be contained. Ca and REM areelements contributing to improvement in stretch flanging propertythrough conformational control of inclusions. If the Ca content exceeds0.1% or the REM content exceeds 0.1%, however, there would be a decreasein cleanliness, and a decrease in ductility.

The balance of the composition of the steel sheet is Fe and incidentalimpurities. Allowable incidental impurities are Sb: not more than 0.01%,Sn: not more than 0.1%, Zn: not more than 0.01%, Co: not more than 0.1%,Zr: not more than 0.1%, and B: not more than 0.1%.

A method for manufacturing the hot-rolled steel sheet of the presentinvention will now be described.

The hot-rolled steel sheet of the present invention is made byhot-rolling a steel slab having a composition within the rangesdescribed above into a prescribed thickness.

While the steel slab used is preferably manufactured by a continuouscasting process to prevent macro-segregation of the constituents, it maybe manufactured by an ingot casting process or a thin-slab castingprocess. A conventional process employed in this embodiment includes thesteps of manufacturing a steel slab, cooling the steel slab to roomtemperature, and reheating the slab. Alternatively, an energy-savingprocess also is applicable without problem in the present invention. Forexample, a hot steel slab, is charged into a heating furnace withoutcooling to room temperature, or directly rolled immediately after shorttemperature holding(direct-hot-charge rolling or direct rolling).

The reheating temperature SRT of the material (steel slab) is notlimited and is preferably not less than 900° C.

Slab reheating temperature: not less than 900° C.

The slab reheating temperature (SRT) is preferably the lowest possiblewith a view to prevent surface defects caused by Cu when the materialcontains Cu. However, with a reheating temperature of less than 900° C.,there is an increase in the rolling load, thus increasing the risk ofoccurrence of a trouble during hot rolling. Considering the increase inscale loss caused along with accelerated oxidation, the slab reheatingtemperature is preferably not more than 1,300° C.

From the viewpoint of reducing the slab reheating temperature andpreventing occurrence of troubles during hot rolling, use of a so-calledsheet bar heater heating a sheet bar is of course an effective method.

The reheated steel slab is then hot-rolled into a hot-rolled sheet. Inthe present invention, a finish rolling condition is particularlyimportant, and the hot rolling is preferably performed at a finishrolling end temperature (FDT) in the range of 780 to 980° C.

At the FDT of less than 780° C., a deformed structure remains in thesteel sheet to cause deterioration of ductility. On the other hand, anFDT exceeding 980° C. coarsens the structure, leading to a decrease informability due to delay of ferrite transformation. Thus, the FDT ispreferably in the range of 780 to 980° C.

After the completion of finish rolling, a forced cooling treatment isapplied. In the present invention, a forced cooling condition isparticularly important. In the present invention, within 2 seconds afterthe completion of finish rolling, a forced cooling is preferably carriedout at a cooling rate of not less than 50° C./second to a temperature inthe range of 620 to 780° C. With a cooling start time exceeding 2seconds, the structure coarsens and ferrite transformation is delayed,resulting in poor press formability. The cooling start time after thecompletion of finish rolling is preferably limited to within 2 seconds.

With a cooling rate of less than 50° C./second after the completion offinish rolling, and ferrite transformation undesirably starts during theforced cooling, ferrite transformation does not appropriately occur in asubsequent isothermal holding treatment or slow cooling treatment, thusresulting in a decreased press formability. Accordingly, the coolingrate is preferably limited to not less than 50° C./second. However, witha cooling rate exceeding 300° C./second, degradation of the steel sheetshape is concerned. Thus, the upper limit of the cooling rate ispreferably 300° C./second.

In addition, in the present invention, the steel sheet is preferablycooled to the vicinity of a nose of a free or pro-eutectoid ferritetemperature region of 620 to 780° C. by the above-mentioned forcedcooling. At a cooling stop temperature of less than 620° C. of theforced cooling, free ferrite is not generated, but pearlite isgenerated. At a cooling stop temperature exceeding 780° C., a decreasein concentration of carbon into austenite decreases with a decrease inthe generation of free ferrite. The cooling stop temperature of forcedcooling is more preferably in the range of 650 to 750° C.

After the forced cooling to the vicinity of a nose of free ferritetemperature region of 620 to 780° C., an isothermal holding treatmentfor 1 to 10 seconds within the above-mentioned temperature region or aslow cooling treatment at a cooling rate of not more than 20° C./secondis preferably performed.

By the isothermal holding treatment for a short period of time withinthis temperature region (620 to 780° C.) or the slow cooling treatmentfor a short period of time within the above-mentioned temperatureregion, a desired amount of free ferrite can be formed.

For achieving the concentration of carbon into the austenite along withferrite transformation, the isothermal holding treatment or slow coolingtreatment is more preferably performed within a temperature region of620° C. to 750° C.

A holding time of the isothermal treatment or a time required for theslow cooling treatment of less than 1 second causes insufficientconcentration of carbon into the austenite. On the other hand, a timeexceeding 10 seconds causes pearlite transformation.

A cooling rate of the slow cooling treatment exceeding 20° C./secondcauses insufficient concentration of carbon into the austenite.

After the isothermal holding treatment or slow cooling treatment, therolled sheet is preferably cooled again to a temperature of 300 to 500°C. at a cooling rate of not less than 50° C./second, and then coiled.That is, the rolled sheet is preferably coiled at a coiling temperature(CT) of 300 to 500° C.

After the isothermal holding treatment or slow cooling treatment, therolled sheet is cooled to a temperature of 300 to 500° C. Also, thecooling rate of this treatment is preferably not less than 50°C./second. With the cooling rate of less than 50° C./second, pearlitetransformation occurs and ductility is decreased. The cooling rate ismore preferably within the range of 50 to 200° C./second.

With a coiling temperature CT of less than 300° C., the secondary phasecontains martensite. On the other hand, with the coiling temperatureexceeding 500° C., the secondary phase contains pearlite. Thus, thecoiling temperature CT is preferably within a range of 300 to 500° C.

In the present invention, all or part of finish rolling may belubrication rolling to reduce the rolling load during hot rolling.Application of lubrication rolling is effective also from the viewpointof achieving a uniform steel sheet shape and uniform material quality.The frictional coefficient on the lubrication rolling is preferably inthe range of 0.25 to 0.10. A continuous rolling process is preferableone,in which neighboring sheet bars can be connected to each other toperform finish rolling continuously. Application of the continuousrolling process is desirable also from the viewpoint of operationalstability of hot rolling.

After the completion of hot rolling, temper rolling of not more than 10%may be applied for adjustment such as shape correction or surfaceroughness control.

The hot-rolled steel sheet of the invention may be used as a steel sheetfor processing and as a steel sheet for surface treatments. Surfacetreatments include galvanizing (including alloying), tin-plating andenameling. After annealing or galvanizing, the hot-rolled steel sheet ofthe present invention may be subjected to a special treatment to improveactivity to chemical treatment, weldability, press formability, andcorrosion resistance.

(2) Cold-rolled Steel Sheet

A cold-rolled steel sheet of the present invention will now bedescribed.

The cold-rolled steel sheet of the present invention has a compositestructure comprising a ferrite primary phase and a secondary phasecontaining retained austenite having a volume ratio of not less than 1%of the entire structure. As described above, a cold-rolled steel sheethaving such a composite structure exhibits high elongation (El), highstrength/elongation balance (TS×El), and excellent press formability.

The volume ratio of the ferrite primary phase contained in the compositestructure is preferably not less than 50%. With a ferrite phase contentof less than 50%, it is difficult to keep high ductility, resulting inpoor press formability. When further enhanced ductility is required, thevolume ratio of the ferrite phase is preferably not less than 80%. Forthe purpose of making full use of advantages of the composite structure,the ferrite phase is preferably not more than 98%.

In the present invention, the steel sheet must contain a retainedaustenite phase as the secondary phase in a volume ratio of not lessthan 1% of the entire structure. With a retained austenite phase contentof less than 1%, it is impossible to obtain high elongation (El). Toobtain higher elongation (El), the retained austenite phase ispreferably contained in a volume ratio of not less than 2%, morepreferably, not less than 3%.

The secondary phase may be a single retained austenite phase having avolume ratio of not less than 1%, or may be a mixture of a retainedaustenite phase of a volume ratio of not less than 1% and an auxiliary(another) phase comprising a pearlite phase, a bainite phase, and/or amartensite phase.

The reasons for limiting the composition of the cold-rolled steel sheetof the present invention will now be described. The weight percent inthe composition will simply be denoted hereinafter as %.

C: not more than 0.20%

C is an element, which improves strength of a steel sheet and promotesthe formation of a composite structure of a ferrite phase and a retainedaustenite phase, and is preferably contained in an amount of not lessthan 0.01% from the viewpoint of forming the retained austenite phase inthe present invention. A C content is more preferably not less than0.05%. A C content exceeding 0.20%, however, causes an increase inamount of carbides in the steel, resulting in a decrease in ductility,and hence a decrease in press formability. A more serious problem isthat a C content exceeding 0.20% leads to remarkable deterioration ofspot weldability and arc weldability. For these reasons, in the presentinvention, the C content is limited to not more than 0.20%. From theviewpoint of formability, the C content is preferably not more than0.18%.

Si: not more than 2.0%

Si is a useful strengthening element, which improves strength of a steelsheet without a marked decrease in ductility of the steel sheet andfacilitates the formation of a residual austenite phase. The Si contentis preferably not less than 0.1%. An Si content exceeding 2.0%, however,leads to deterioration of press formability and degrades the surfacequality. The Si content is, therefore, limited to not more than 2.0%.

Mn: not more than 3.0%

Mn is a useful element, which strengthens the steel and prevents hotcracking caused by S, and is therefore contained in an amount accordingto the S content. These effects are particularly remarkable at an Mncontent of not less than 0.5%. However, an Mn content exceeding 3.0%results in deterioration of press formability and weldability. The Mncontent is, therefore, limited to not more than 3.0% in the presentinvention. More preferably, the Mn content is not less than 1.0%.

P: not more than 0.10%

P strengthens the steel, and may be contained in an amount of preferablynot less than 0.005%, according to a desired strength. However, anexcess P content causes deterioration of press formability. The Pcontent is, therefore, limited to not more than 0.10%. When moreexcellent press formability is required, the P content is preferably notmore than 0.08%.

S: not more than 0.02%

S is an element, which is present as inclusions in steel and causesdeterioration of ductility, formability, and particularly stretchflanging formability of a steel sheet, and it should be the lowestpossible. However, an S content reduced to not more than 0.02% does notexert much adverse effect. Thus, the S content is limited to not morethan 0.02% in the present invention. When superior stretch flangingformability is required, the S content is preferably not more than0.010%.

Al: not more than 0.30%

Al is a deoxidizing element of steel, and is useful for improvingcleanliness of the steel. In addition, Al is effective for the formationof the retained austenite. In order to obtain these effects, the Alcontent is preferably not less than 0.01%. However, an Al contentexceeding 0.30% cannot give further enhanced deoxidizing effects, andcauses deterioration of press formability. The Al content is, therefore,limited to not more than 0.30%. The invention also includes a steelmaking process using other deoxidizers, for example, Ti or Si, and steelsheets produced by such deoxidation methods are also included in thescope of the invention. In this case, addition of Ca or REM to moltensteel does not impair the features of the steel sheet of the inventionat all. Of course, steel sheets containing Ca or REM are included withinthe scope of the invention.

N: not more than 0.02%

N is an element, which increases strength of a steel sheet through solidsolution strengthening or strain age hardening, and is preferablycontained in an amount of not more than 0.001%. However, an N contentexceeding 0.02% causes an increase in nitride content in the steelsheet, whereby ductility and press formability of the steel sheet areseriously deteriorated. The N content is therefore limited to not morethan 0.02%. When further improvement of press formability is required,the N content is preferably not more than 0.01%.

Cu: 0.5 to 3.0%

Cu is an element, which remarkably increases strain age hardening of asteel sheet (increase in strength after pre-deformation/heat treatment),and is one of the most important elements in the present invention. Witha Cu content of less than 0.5%, an increase in tensile strength ΔTSexceeding 80 MPa cannot be obtained by changing the pre-deformation/heattreatment conditions. In the present invention, therefore, Cu should becontained in an amount of not less than 0.5%. With a Cu contentexceeding 3.0%, however, the effect is saturated, leading to unfavorableeconomic effects. Furthermore, deterioration of press formabilityoccurs, and the surface quality of the steel sheet is degraded. The Cucontent is, therefore, limited within the range of 0.5 to 3.0%. In orderto simultaneously achieve a higher ΔTS and excellent press formability,the Cu content is preferably within the range of 1.0 to 2.5%.

In the present invention, the above-mentioned composition containing Cupreferably further contains, in weight percent, at least one of thefollowing Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

Group A: Ni: not more than 2.0%

Group A: Ni is an element effective for preventing surface defectsproduced by Cu contained in the steel sheet, and may be contained asrequired. The Ni content depends on the Cu content, and is preferablyabout a half the Cu content, more specifically, within the range ofabout 30 to about 80% of the Cu content. An Ni content exceeding 2.0%cannot give further enhancement in the effect because of saturation ofthe effect, leading to economic disadvantages, and causes deteriorationof press formability. For these reasons, the Ni content is preferablylimited to not more than 2.0%.

Group B: at least one of Cr and Mo: not more than 2.0% in total

Group B: Both Cr and Mo, as well as Mn, strengthen the steel sheet andmay be contained as required preferably in an amount of not less than0.1% for Cr and not less than 0.1% for Mo. If at least one of Cr and Moare contained in an amount exceeding 2.0% in total, press formability isimpaired. It is therefore preferable to limit the total content of Crand Mo forming Group B to not more than 2.0%.

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total

Group C: Nb, Ti, and V are elements, which effectively form finedispersion of carbides contributing to an increase in strength.Therefore, Nb, Ti, and V can be selected and contained as requiredpreferably in an amount of not less than 0.01% for Nb, in an amount ofnot less than 0.01% for Ti and in an amount of not less than 0.01% forV. If the total content of at least one of Nb, Ti, and V exceeds 0.2%,the press formability is impaired. Thus, the total content of Nb, Tiand/or V is preferably limited to not more than 0.2%.

In the present invention, in place of the aforementioned Cu, at leastone selected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to2.0%, and W: 0.05 to 2.0% may be contained in an amount of not more than2.0% in total.

At least one selected from the group consisting of Mo: 0.05 to 2.0%, Cr:0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not more than 2.0% intotal

In the present invention, all of Mo, Cr, and W, as well as Cu, are themost important elements, which remarkably increase strain age hardeningof the steel sheet, and can be selected and contained. When a steelsheet containing at least one of Mo, Cr, and W and having a compositestructure of a ferrite phase and a phase containing retained austeniteis subjected to a prestrain (pre-deformation) of not less than 5% and alow-temperature heat treatment (heat treatment), the retained austeniteis changed into martensite by strain-induced transformation. Then, theformation of fine carbide precipitation in the martensite is induced bythe strain, resulting in an increase in tensile strength ΔTS of not lessthan 80 MPa. With a content of each of these elements of less than0.05%, changing pre-deformation/heat treatment conditions does not givean increase in tensile strength ΔTS of at least 80 MPa. If the contentof each of these elements exceeds 2.0%, a further enhanced effectcorresponding to the content cannot be expected as a result ofsaturation of the effect, leading to economic disadvantages, and thisresults in deterioration of press formability. The contents of Mo, Cr,and W are therefore limited within the range of 0.05 to 2.0% for Mo,0.05 to 2.0% for Cr, and 0.05 to 2.0% for W. From the viewpoint of pressformability, the total content of Mo, Cr, and W is limited to not morethan 2.0%.

In the present invention, at least one selected from the groupconsisting of Mo, Cr, and W is preferably contained and further, atleast one of Nb, Ti, and V are preferably contained not more than 2.0%in total.

At least one of Nb, Ti, and V, in a total amount of not more than 2.0%:

Nb, Ti, and V are elements forming carbides, and can be selected andcontained as required, when at least one of Mo, Cr, and W is added. Whenthe steel composition contains at least one of Mo, Cr, and W and has acomposite structure containing a ferrite phase and a retained austenitephase, and contains at least one of Nb, Ti, and V, the retainedaustenite is transformed into martensite by strain-inducedtransformation during the pre-deformation/heat treatment. Then, finecarbide precipitation is induced by the strain in the martensite, thusresulting in an increase in tensile strength ΔTS of not less than 80MPa. This effect is particularly remarkable preferably at a Nb contentof not less than 0.01%, at a Ti content of not less than 0.01%, and at aV content of not less than 0.01%. However, a total content of Nb, Ti,and V exceeding 2.0% causes deterioration of press formability. Thus,the total content of Nb, Ti and/or V is preferably limited to not morethan 2.0%.

Although no particular restriction is imposed, apart from theabove-mentioned constituents, the composition may contain B: not morethan 0.1%, Zr: not more than 0.1%, Ca: not more than 0.1%, and REM: notmore than 0.1% without any problem.

The balance of the composition of the steel is Fe and incidentalimpurities. Allowable incidental impurities include Sb: not more than0.01%, Sn: not more than 0.1%, Zn: not more than 0.01%, and Co: not morethan 0.1%.

The method for manufacturing the cold-rolled steel sheet of the presentinvention will now be described.

The cold-rolled steel sheet of the present invention is manufacturedthrough a hot rolling step of hot-rolling a steel slab having thecomposition within the aforementioned ranges into a hot-rolled steelsheet, a cold rolling step, of cold-rolling the hot-rolled steel sheetinto a cold-rolled steel sheet, and a recrystallization annealing stepof recrystallization-annealing the cold-rolled steel sheet to form acold-rolled annealed steel sheet.

Although the steel slab used is preferably manufactured by a continuouscasting process to prevent macrosegregation of the constituents, it maybe manufactured by an ingot casting process or a thin-slab continuouscasting process. A conventional process employed in this embodimentincludes the steps of manufacturing a steel slab, cooling the steel slabto room temperature, and reheating the slab. Alternatively, anenergy-saving process is applicable without problem in the presentinvention. For example, a hot steel slab is charged into a reheatingfurnace without cooling to room temperature, or directly rolledimmediately after short temperature holding (direct-feed rolling ordirect rolling).

The steel slab having the above-mentioned composition is reheated andhot-rolled to make a hot-rolled steel sheet. No particular problem isencountered as to conventionally known conditions so far as suchconditions permit manufacture of a hot-rolled steel sheet having adesired thickness in the hot rolling step. Preferable conditions for hotrolling are as follows:

Slab reheating temperature: not less than 900° C.

The slab reheating temperature is preferably the lowest possible with aview to prevent surface defects caused by Cu when the compositioncontains Cu. However, with a reheating temperature of less than 900° C.,the rolling load increases, thus increasing the risk of occurrence of atrouble during hot rolling. In view of an increase in scale loss causedby facilitated oxidation, the slab reheating temperature is preferablynot more than 1,300° C.

From the viewpoint of reducing the slab reheating temperature andpreventing occurrence of troubles during hot rolling, use of a so-calledsheet bar heater, which heats a sheet bar, is effective.

Finish rolling end temperature: not less than 700° C.

At a finish rolling end temperature (FDT) of not less than 700° C., itis possible to obtain a uniform hot-rolled mother sheet structure whichcan give an excellent formability after cold rolling andrecrystallization annealing. A finish rolling end temperature of lessthan 700° C. leads to a non-uniform structure of the hot-rolled mothersheet and a higher rolling load during hot rolling, thus increasing therisk of occurrence of troubles during hot rolling. Thus, the FDT for thehot rolling step is preferably not less than 700° C.

Coiling temperature: not more than 800° C.

The coiling temperature is preferably not more than 800° C., and morepreferably not less than 200° C. A coiling temperature exceeding 800° C.tends to cause a decrease in yield as a result of an increased scaleloss. With a coiling temperature of less than 200° C., the steel sheetshape is seriously impaired, and there is an increasing risk ofoccurrence of inconveniences in practical use.

In the hot rolling step in the present invention, as described above, itis desirable to reheat the slab to a temperature of not less than 900°C., hot-roll the reheated slab at a finish rolling end temperature ofnot less than 700° C., and coil the hot-rolled steel sheet at a coilingtemperature of not more than 800° C. and preferably not less than 200°C.

In the hot rolling step in the present invention, all or part of finishrolling may be lubrication rolling, which reduces the rolling loadduring the hot rolling. The lubrication rolling is effective also fromthe viewpoint of achieving a uniform steel sheet shape and a uniformmaterial quality. The frictional coefficient on the lubrication rollingis preferably within a range of 0.25 to 0.10. It is desirable to connectneighboring sheet bars to each other to perform a continuous finishrolling process. Application of the continuous rolling process isdesirable also from the viewpoint of operational stability of hotrolling.

Then, a cold rolling step is conducted for the hot-rolled steel sheet.In the cold rolling step, the hot-rolled steel sheet is cold-rolled intoa cold-rolled steel sheet. Any cold rolling conditions may be used sofar as such conditions permit production of cold-rolled steel sheetswith desired dimensions and shape, and no particular restriction isimposed. The reduction in cold rolling is preferably not less than 40%.With a reduction of less than 40%, uniform recrystallization barelyoccurs during the subsequent recrystallization-annealing step.

Then, the cold-rolled steel sheet is subjected to the recrystallizationannealing step to convert the sheet into a cold-rolled annealed steelsheet. The recrystallization annealing is preferably carried out on acontinuous annealing line. In the present invention, therecrystallization annealing is a heat treatment which includes heatingand soaking the cold-rolled sheet in the dual phase region of ferriteand austenite in the temperature range between the A_(C1) transformationpoint and the A_(C3) transformation point, cooling the sheet, andretaining the sheet at a temperature in the range of 300 to 500° C. for30 to 1,200 seconds.

The heating and soaking temperature for recrystallization annealing ispreferably within the dual phase region in the temperature range betweenthe A_(C1) transformation point and the A_(C3) transformation point. Theheating and soaking temperature of less than the A_(C1) transformationpoint leads to the formation a single ferrite phase. On the other hand,a high temperature exceeding A_(C3) transformation point results incoarsening of crystal grains, the formation of a single austenite phase,and a serious deterioration of press formability.

After the heating and soaking treatment, the sheet was cooled from theheating and soaking temperature and retained at a temperature in therange of 300 to 500° C. for 30 to 1,200 seconds. The heating and soakingtreatment and the subsequent retaining treatment facilitates theformation of a retained austenite phase of not less than 1%. When thetemperature for the retaining treatment is less than 300° C., thecomposite structure of ferrite and martensite is formed. On the otherhand, a temperature range exceeding 500° C. leads to a ferrite/bainitecomposite structure or a ferrite/pearlite composite structure. In thesecases, the retained austenite is barely formed.

In addition, a retention time of less than 30 seconds in the temperaturerange of 300 to 500° C. cannot lead to the formation of the retainedaustenite structure. Also, the retention time exceeding 1,200 secondscannot lead to the formation of the retained austenite structure, butleads to a ferrite/bainite composite structure. Therefore, the retentiontime in the temperature region of 300 to 500° C. is preferably in therange of 30 to 1,200 seconds.

By the recrystallization annealing, a composite structure of a ferritephase and a retained austenite phase is formed, whereby a high ΔTS canbe obtained together with high ductility.

After the hot rolling, temper rolling with a reduction rate of not morethan 10% may be applied for adjustments and other shape correction and,surface roughness control.

The cold-rolled steel sheet of the invention may be used as a steelsheet for processing and as a steel sheet for surface-treating. Surfacetreatments include galvanizing (including alloying), tin-plating andenameling. After galvanizing, the cold-rolled steel sheet of the presentinvention may be subjected to a special treatment to improve activity tochemical treatment, weldability, press formability, and corrosionresistance.

(3) Hot-dip Galvanized Steel Sheet

The hot-dip galvanized steel sheet of the present invention will now bedescribed.

The hot-dip galvanized steel sheet of the present invention has acomposite structure comprising a primary phase consisting of a ferritephase and a tempered martensite phase and a secondary phase containingretained austenite phase in a volume ratio of not less than 2%.

Note that the term “tempered martensite phase” in the present inventionmeans a phase produced by heating a lath martensite. That is, thetempered martensite phase still maintains a fine internal structure ofthe lath martensite, after the heating (tempering). Furthermore, thetempered martensite phase is softened by heating (tempering), hasenhanced deformability as compared with martensite, and is effective forimproving ductility of the steel sheet. Note that the term “lathmartensite” means martensite consisting of bundles of thin longplatelike martensite crystals, which can be observed with an electronmicroscope.

In the hot-dip galvanized steel sheet of the present invention, thetotal volume ratio of the ferrite phase and the tempered martensitephase functioning as the primary phase is preferably not less than 50%.With a total volume ratio of the ferrite phase and the tempered phase ofless than 50%, it is difficult to secure high ductility and pressformability is decreased. When further enhanced ductility is required,the total volume ratio of the ferrite phase and the martensite phasefunctioning as the primary phase is preferably not less than 80%. Forthe purpose of making full use of advantages of the composite structure,the total of the ferrite phase and the tempered martensite phase ispreferably not more than 98%. The ferrite phase constituting the primaryphase preferably occupies not less than 30% by volume of the entirestructure, and the tempered martensite phase preferably occupies notless than 20% by volume of the entire structure. With a volume ratio ofthe ferrite phase of less than 30%, or with a volume ratio of thetempered martensite phase of less than 20%, the ductility will not beremarkably enhanced.

The hot-dip galvanized steel sheet of the present invention contains aretained austenite phase as a secondary phase with a volume ratio of notless than 1% of the entire structure. With a content of the retainedaustenite phase of less than 1%, high elongation (El) cannot beobtained. In order to obtain higher elongation (El), the retainedaustenite phase is preferably contained not less than 2% and morepreferably not less than 3%. The secondary phase may be a singleretained austenite phase having a volume ratio of not less than 1%, ormay be a mixture of a retained austenite phase of a volume ratio of notless than 1% and an auxiliary (other) phase, for example, a pearlitephase, a bainite phase, and/or a martensite phase.

The reasons for limiting the composition of the hot-dip galvanized steelsheet of the present invention will now be described.

C: not more than 0.20%

C is an element, which improves the strength of a steel sheet andpromotes the formation of a composite structure of a primary phasecomprising ferrite and tempered martensite and a secondary phasecontaining retained austenite. In the present invention, from theviewpoint of formation of the composite structure, C is preferablycontained in an amount of not less than 0.01%. A C content exceeding0.20% causes an increase in carbide content in the steel, resulting in adecrease in ductility, and hence a decrease in press formability. A moreserious problem is that a C content exceeding 0.20% leads to remarkabledeterioration of spot weldability and arc weldability. For thesereasons, in the present invention, the C content is limited to not morethan 0.20%. From the viewpoint of formability, the C content ispreferably not more than 0.18%.

Si: not more than 2.0%

Si is a useful strengthening element, which improves strength of a steelsheet without a marked decrease in ductility of the steel sheet, and isnecessary for obtaining retained austenite. These effects areparticularly remarkable at an Si content of not less than 0.1% andtherefore, the Si content is preferably not less than 0.1%. An Sicontent exceeding 2.0%, however, leads to deterioration of pressformability and degrades the platability. Therefore, the Si content islimited to not more than 2.0%.

Mn: not more than 3.0%

Mn is a useful element, which strengthens the steel and prevents hotcracking caused by S, and is therefore contained in an amount accordingto S content. These effects are particularly remarkable at an Mn contentof not less than 0.5%. However, an Mn content exceeding 3.0% results indeterioration of press formability and weldability. The Mn content is,therefore, limited to not more than 3.0%. More preferably, the Mncontent is not less than 1.0%.

P: not more than 0.10%

P strengthens the steel. In the present invention, P is preferablycontained in an amount of not less than 0.005% for securing thestrength. However, an excess content of P exceeding 0.10% causesdeterioration of press formability. For this reason, in the presentinvention, a P content is limited to not more than 0.10%. When moreenhanced press formability is required, the P content is preferably notmore than 0.08%.

S.: not more than 0.02%

S is an element, which is present as inclusions in a steel sheet andcauses deterioration of ductility, formability, and particularly stretchflanging formability of the steel sheet, and it should be the lowestpossible. An S content reduced to not more than 0.02% does not exertmuch adverse effect and therefore, the S content is limited to not morethan 0.02% in the present invention. When excellent stretch flangingformability is required, the S content is preferably not more than0.010%.

Al: not more than 0.10%

Al is a deoxidizing element of steel, and is useful for improvingcleanliness of steel. In addition, Al is effective for the formation ofthe retained austenite. In the present invention, the Al content ispreferably not less than 0.01%. An excess Al content exceeding 0.30%,however, cannot give a further enhanced effect because of saturation ofthe effect, and causes deterioration of press formability. The Alcontent is, therefore, limited to not more than 0.30%. The presentinvention also include a steel making process using other deoxidizers,for example, Ti or Si, and steel sheets produced by such deoxidationmethods are also included in the scope of the present invention. In thiscase, addition of Ca or REM to molten steel does not impair the featuresof the steel sheet of the present invention at all. Of course, steelsheets containing Ca or REM are included within the scope of the presentinvention.

N: not more than 0.02%

N is an element, which increases strength of a steel sheet through solidsolution strengthening or strain age hardening, and is preferablycontained in an amount of not less than 0.001%. An N content exceeding0.02% causes an increase in the nitride content in the steel sheet,which causes serious deterioration of ductility and of pressformability. The N content is, therefore, limited to not more than0.02%. When further improvement of press formability is required, the Ncontent is preferably not more than 0.01%.

Cu: 0.5 to 3.0%

Cu is an element, which remarkably increases strain age hardening of asteel sheet (increase in strength after pre-deformation/heat treatment),and is the most important element in the present invention. With a Cucontent of less than 0.5%, an increase in tensile strength ΔTS of notless than 80 MPa cannot be obtained by changing the pre-deformation/heattreatment conditions. In the present invention, therefore, Cu should becontained in an amount of not less than 0.5%. With a Cu contentexceeding 3.0%, however, the effect is saturated, leading to unfavorableeconomic effects. Furthermore, deterioration of press formabilityoccurs, and the surface quality of the steel sheet is degraded. The Cucontent is, therefore, limited within the range of 0.5 to 3.0%. In orderto simultaneously achieve a higher ΔTS and excellent press formability,the Cu content is preferably within the range of 1.0 to 2.5%.

In the present invention, it is preferable that the compositioncontaining Cu further contain, in weight percent, at least one of thefollowing Groups A to C:

Group A: Ni: not more than 2.0%;

Group B: at least one of Cr and Mo: not more than 2.0% in total; and

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total.

Group A: Ni: not more than 2.0%

Group A: Ni is an element effective for preventing surface defectsproduced by Cu contained in the steel sheet, and can be contained asrequired. The Ni content depends on the Cu content, and is preferablyabout a half the Cu content, more specifically, within the range ofabout 30 to about 80% of the Cu content. An Ni content exceeding 2.0%cannot give further enhancement in the effect because of saturation ofthe effect, leading to economic disadvantages, and causes deteriorationof press formability. For these reasons, the Ni content is preferablylimited to not more than 2.0%.

Group B: at least one of Cr and Mo: not more than.2.0% in total

Group B: Both Cr and Mo strengthen the steel sheet, like Mn, and can becontained as required. However, if at least one of Cr and Mo arecontained in an amount exceeding 2.0% in total, press formability isimpaired. The total content of Cr and Mo is preferably limited to notmore than 2.0%. From the viewpoint of press formability, a Cr content ispreferably not less than 0.1%, and an Mo content is preferably not lessthan 0.1%.

Group C: at least one of Nb, Ti, and V: not more than 0.2% in total

Group C: Nb, Ti, and V are carbide-forming elements and increasestrength by fine dispersion of carbides, and can be selected andcontained as required. However, if the total content of at least one ofNb, Ti, and V exceeds 0.2%, press formability is impaired. Thus, thetotal content of Nb, Ti and V is preferably limited to not more than0.2%. The above-mentioned effect can be achieved at an Nb content of notless than 0.01%, at a Ti content of not less than 0.01%, and at a Vcontent of not less than 0.01%.

In the present invention, in place of Cu, at least one selected from thegroup consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0%, Cr, and W: 0.05to 2.0% may be contained in an amount of not more than 2.0% in total.

At least one selected from the group consisting of Mo: 0.05 to 2.0%, Cr:0.05 to 2.0% and W: 0.05 to 2.0%, in an amount of not more than 2.0% intotal

In the present invention, all of Mo, Cr, and W, as well as Cu, are themost important elements, which remarkably increase strain age hardening(increase in strength after pre-deformation/heat treatment) of the steelsheet. When a steel sheet containing at least one of Mo, Cr, and W, andhaving a composite structure comprising a primary phase of a ferritephase and a tempered martensite phase and a secondary phase containingretained austenite in a volume ratio of not less than 1% is subjected toprestrain (pre-deformation) of not less than 5% and a low-temperatureheat treatment (heat treatment), the retained austenite is transformedinto martensite by strain-induced transformation. Then, the formation offine carbide precipitation is induced by the strain at a low temperatureoccurs in the martensite, resulting in an increase in tensile strengthΔTS of not less than 80 MPa. With a content of each of these elements ofless than 0.05%, changing the steel sheet structure andpre-deformation/heat treatment conditions does not give an increase intensile strength ΔTS of not less than 80 MPa. Therefore, in the presentinvention, each of Mo, Cr, and W is preferably contained in an amount ofnot less than 0.05%. If the content of each of Mo, Cr, and W eachexceeds 2.0%, a further enhanced effect corresponding to the contentcannot be expected as a result of saturation of the effect, leading toeconomic disadvantages, and this results in deterioration of pressformability. For these reasons, the content of each of Mo, Cr, and W ispreferably limited within the range of 0.05 to 2.0%, and the totalcontent thereof is preferably limited to not more than 2.0%.

The above-mentioned composition containing at least one of Mo, Cr, and Wpreferably further contains at least one of Nb, Ti, and V in an amountof not more than 2.0% in total.

At least one of Nb, Ti, and V, in a total amount of not more than 2.0%

Nb, Ti, and V are carbide-forming elements and can be selected andcontained as required, when at least one of Mo, Cr, and W is added.However, a total content of Nb, Ti, and V exceeding 2.0% causesdeterioration of press formability. Thus, the total content of Nb, Ti,and V is preferably limited to not more than 2.0%. At least one of Mo,Cr, and W are added, at least one of Nb, Ti, and V are added, and thestructure is transformed into a composite structure of a primary phasecomprising a ferrite phase and a tempered martensite phase and asecondary phase containing retained austenite. This forms fine compositecarbides in martensite which was formed by strain-induced transformationduring the pre-deformation/heat treatment, and strain-induced fineprecipitation at a low temperature occurs, resulting in an increase intensile strength ΔTS of not less than 80 MPa. In order to obtain thiseffect, Nb, Ti, and V is preferably contained in an amount of not lessthan 0.01% for Nb, in an amount of not less than 0.01% for Ti and in anamount of not less than 0.01% for V, and at least one of Nb, Ti, and Vcan be selected and contained as required.

Although no particular restriction is imposed, apart from theabove-mentioned constituents, the composition may contain B: not morethan 0.1%, Ca: not more than 0.1%, Zn: not more than 0.1%, and REM: notmore than 0.1% without any problem.

The balance of the composition of the steel is Fe and incidentalimpurities. Allowable incidental impurities include Sb: not more than0.01%, Sn: not more than 0.1%, Zn: not more than 0.01%, and Co: not morethan 0.1%.

The method for manufacturing the hot-dip galvanized steel sheet of thepresent invention will now be described.

The hot-dip galvanized steel sheet is preferably manufactured through aprimary heat treatment step of heating a steel sheet having theabove-mentioned composition to a temperature of not less than the A_(C1)transformation point and rapidly cooling the steel sheet, a secondaryheat treatment step of heating the steel sheet to a temperature offerrite/austenite dual phase within the range of A_(C1) transformationpoint to A_(C3) transformation point on a continuous hot-dip galvanizingline, and a hot-dip galvanizing step of forming a hot-dip galvanizinglayer on each surface of the steel sheet.

A hot-rolled steel sheet or a cold-rolled steel sheet may preferably beused in this process. A preferable manufacturing method of the steelsheet used will now be described, although the method is not limitedthereto in the present invention.

A suitable method for manufacturing the hot-rolled steel sheet used as agalvanizing substrate will be described.

A material (steel slab) used is preferably manufactured by a continuouscasting process to prevent macro-segregation of the constituents, but itmay be manufactured by an ingot casting process or a thin-slab castingprocess. A conventional process employed in this embodiment includes thesteps of manufacturing a steel slab, cooling the steel slab to roomtemperature, and reheating the slab. Alternatively, an energy-savingprocess is applicable with no problem. As the energy-saving process, forexample, a direct-hot charge rolling process of charging the hot steelslab into a reheating furnace without cooling the same, and a directrolling process of immediately rolling after a short temperature holdingare applicable.

The material (steel slab) is first heated, and subjected to a hotrolling step to form a hot-rolled steel sheet. Known hot rollingconditions may be employed without problem as long as a hot-rolled steelsheet having a desired thickness is formed. Preferable conditions forhot rolling are as follows:

Slab reheating temperature: not less than 900° C.

In the case of a steel slab containing Cu, the slab heating temperatureis preferably the lowest possible to prevent surface defects caused byCu. However, a heating temperature of less than 900° C. causes anincrease in the rolling load, thus increasing the risk of occurrence ofa trouble during the hot rolling. Considering the increase in scale losscaused by accelerated oxidation, the slab heating temperature ispreferably not more than 1,300° C. From the viewpoint of decreasing theslab heating temperature and preventing occurrence of troubles duringhot rolling, use of a so-called sheet bar heater, which heats a sheetbar, is effective.

Finish rolling end temperature: not less than 700° C.

At a finish rolling end temperature FDT of not less than 700° C., it ispossible to obtain a uniform hot-rolled mother sheet structure which cangive an excellent formability after cold rolling and recrystallizationannealing. A finish rolling end temperature FDT of less than 700° C.leads to a non-uniform structure of the hot-rolled mother sheet and ahigher rolling load during hot rolling, thus increasing the risk ofoccurrence of troubles during hot rolling. Thus, the FDT for the hotrolling step is preferably not less than 700° C.

Coiling temperature: not more than 800° C.

The coiling temperature CT is preferably not more than 800° C., and morepreferably not less than 200° C. The CT exceeding 800° C. tends to causea decrease in yield as a result of an increased scale loss. With a CT ofless than 200° C., the steel sheet shape is seriously impaired, andthere is an increasing risk of occurrence of inconveniences in practicaluse.

The hot-rolled steel sheet suitably applicable in the present inventionis preferably prepared by heating the slab to not less than 900° C.,hot-rolling the heated slab at a finish rolling end temperature of notless than 700° C., and coiling the hot-rolled sheet at a coilingtemperature of not less than 800° C., and preferably not less than 200°C.

In the above-mentioned hot rolling step, all or part of finish rollingmay be lubrication rolling, which reduces the rolling load during thehot rolling. The lubrication rolling is effective also from theviewpoint of achieving a uniform steel sheet shape and a uniformmaterial quality. The frictional coefficient on the lubrication rollingis preferably within the range of 0.25 to 0.10. It is desirable toconnect neighboring sheet bars to each other to perform a continuousfinish rolling process. Application of the continuous rolling process isdesirable also from the viewpoint of operational stability of hotrolling.

The hot-rolled sheet with scales may be annealed to form an internaloxide layer at the surface of the steel sheet. The internal oxide layer,which prevents concentration of Si, Mn, and P at the surface, improveshot-dip galvanizing ability.

The hot-rolled sheet manufactured by the above-mentioned method may beused as an original sheet for plating. Alternatively, the hot-rolledsheet may be cold-rolled to form a cold-rolled sheet used as an originalsheet for plating.

In the cold rolling step, any cold rolling condition may be used withoutparticular restriction so far as such a condition permits production ofcold-rolled steel sheets with desired dimensions and shapes. Thereduction in cold rolling is preferably not less than 40%. A reductionof less than 40% inhibits uniform recrystallization during thesubsequent primary heat treatment.

In the present invention, the above-mentioned steel sheet (hot-rolledsheet or cold-rolled sheet) is subjected to a primary heat treatmentstep including heating to a temperature of not less than the A_(C1)transformation point and rapid cooling.

Heating in the primary heat treatment, the steel sheet is preferablyheld at a temperature of not less than A_(C1) transformation point, morepreferably not less than (A_(C3) transformation point—50° C.), and mostpreferably not less than A_(C3) transformation point. After heating, thesteel sheet is preferably rapidly cooled to a temperature of not morethan the Ms point at a cooling rate of not less than 10° C./second.During the primary heat treatment step, lath martensite is produced inthe steel sheet. In the present invention, the most important point isto form lath martensite during the primary heat treatment step. Unlessthe lath martensite is formed in the steel sheet, it is difficult toform a secondary phase containing retained austenite in the subsequentsteps.

When a hot-rolled steel sheet, subjected to final hot rolling at atemperature of not less than (Ar₃ transformation point−50° C.), is usedas an original sheet for plating, the primary heat treatment step can besubstituted the steel sheet for rapidly cooling to a temperature of notless than Ms point at a cooling rate of not less than 10° C./secondduring cooling after the final hot rolling.

Then, the steel sheet containing lath martensite formed during theabove-described primary heat treatment is subjected to a secondary heattreatment step for heating to and holding at a temperature in the rangeof A_(C1) transformation point to A_(C3) transformation point on acontinuous hot-dip galvanizing line. During the secondary heat treatmentstep, the lath martensite formed during the primary heat treatment stepis changed into tempered martensite, and a part of the structure istransformed into austenite for formation of retained austenite.

A heating and holding temperature of less than the A_(C1) transformationpoint in the secondary heat treatment step cannot form retainedaustenite. A heating and holding temperature exceeding the A_(C3)transformation point causes retransformation of the entire structure ofthe steel sheet to austenite, whereby the tempered martensitedisappears. For these reasons, the heating and holding temperature inthe secondary heat treatment is within the range of the A_(C1)transformation point to the A_(C3) transformation point.

Then, the steel sheet heated to and held at a temperature in the rangeof the A_(C3) transformation point to the A_(C3) transformation point inthe second heat treatment step is preferably cooled to a temperature ofnot more than 500° C. at a cooling rate of 5° C./second or more, fromthe viewpoint of forming retained austenite. This can achieve acomposite structure of a primary phase containing a ferrite phase and atempered martensite phase and a secondary phase containing retainedaustenite in the steel sheet.

The steel sheet after the secondary heat treatment is subsequentlysubjected to a hot-dip galvanizing treatment step on a continuoushot-dip galvanizing line.

The hot-dip galvanizing treatment may be carried out under treatmentconditions (galvanizing bath temperature: 450 to 500° C.) used in ausual continuous hot-dip galvanizing line without a particularrestriction. Because galvanizing at an excessively high temperatureleads to a poor platability, galvanizing is preferably conducted at atemperature of not more than 500° C. Galvanizing at a temperature ofless than 450° C. causes deterioration of platability. From theviewpoint of forming martensite, the cooling rate from the hot-dipgalvanizing temperature to 300° C. is preferably not less than 5°C./second.

For the purpose of adjusting the galvanizing weight as required aftergalvanizing, wiping may be performed.

After the hot-dip galvanizing treatment, an alloying treatment of agalvanizing layer may be applied. The alloying treatment is preferablycarried out by reheating the plated sheet to a temperature in the rangeof 450 to 500° C. after the hot-dip galvanizing treatment. At analloying treatment temperature of less than 450° C, alloying isdecelerated, resulting in low productivity. On the other hand, analloying treatment temperature exceeding 550° C. causes deterioration ofplatability, makes it difficult to secure a required amount of retainedaustenite, and decrease ductility of the steel sheet.

After the alloying treatment, the sheet is preferably cooled to 300° C.at a cooling rate of not less than 5° C./second. An extremely lowcooling rate after the alloying treatment makes it difficult to form arequired amount of retained austenite.

In the present invention, pickling treatment for removing a concentratedsurface layer of the constituents formed on the surface of the steelsheet during the primary heat treatment step is preferably performedbetween the primary heat treatment step and the hot-dip galvanizingstep, for the improvement in platability. By the primary heat treatment,P and oxides of Si, Mn, Cr, etc. are concentrated on the steel surfaceto form a concentrated surface layer. It is favorable for improvingplatability to remove this concentrated surface layer through picklingand to conduct annealing in a reducing atmosphere subsequently on thecontinuous hot-dip galvanizing line.

After the hot-dip galvanizing or the alloying treatment step, a temperrolling step with a reduction of not more than 10% may be applied foradjustments such as shape correction and surface roughness adjustment.

To the steel sheet of the present invention, any special treatment maybe applied after the hot-dip galvanizing, to improve chemical treatmentability, weldability, press formability, and corrosion resistance.

EXAMPLES Example 1

Molten steels having the compositions shown in Table 1 were made in aconverter and cast into steel slabs by a continuous casting process.Each of these steel slabs was reheated, and hot-rolled under conditionsshown in Table 2 into a hot-rolled steel strip (hot-rolled sheet) havinga thickness of 2.0 mm. The hot-rolled sheet was temper-rolled at areduction of 1.0%. TABLE 1 STEEL COMPOSITION (wt. %) NO. C Si Mn P S AlN Cu Ni Cr, Mo, Nb, Ti, V A 0.09 1.45 1.05 0.01 0.003 0.034 0.002 1.52 —— — B 0.12 1.50 1.20 0.01 0.002 0.030 0.002 1.43 0.65 Mo: 0.32 — C 0.101.48 1.35 0.01 0.002 0.028 0.002 1.25 0.52 Cr: 0.53 — D 0.15 1.53 1.450.01 0.003 0.033 0.002 1.33 0.44 — Nb: 0.01, Ti: 0.01, V: 0.01 E 0.121.48 1.55 0.01 0.005 0.032 0.002 0.15 — — — F 0.11 1.50 1.08 0.01 0.0040.032 0.002 0.68 — — — G 0.13 1.52 1.22 0.01 0.004 0.032 0.002 0.98 — —— H 0.12 1.42 1.22 0.01 0.003 0.033 0.002 1.55 0.62 — — I 0.11 1.52 1.520.01 0.003 0.031 0.002 1.49 — Cr: 0.15, — Mo: 0.12 J 0.13 1.43 1.48 0.010.003 0.028 0.002 1.43 — Mo: 0.21 — K 0.15 1.58 1.05 0.01 0.003 0.0300.002 1.52 — — Nb: 0.01 L 0.14 1.60 1.21 0.01 0.003 0.028 0.002 1.48 —Cr: 0.11 Ti: 0.01

TABLE 2 HOT ROLLING - COOLING AFTER ROLLING TIME FINISH BE- FORCED SLOWSLAB ROLLING FORE COOLING ISOTHERMAL COOLING TREATMENT COOLING COIL-REHEATING END START COOL- HOLDING COOL- RATE ING STEEL TEMP. TEMP. COOL-ING STOP HOLD- INITIAL ING STOP BEFORE TEMP. SHEET STEEL SRT FDT INGRATE TEMP. TEMP. ING TEMP. RATE TEMP. COILING CT NO. NO. ° C. ° C. S °C./s ° C. ° C. TIME S ° C. ° C./s ° C. ° C./s ° C. 1 A 1250 850 0.5 100710 710 5 — — — 80 450 2 B 1250 850 0.5 80 690 690 5 — — — 60 450 3 1250850 0.3 30 700 — — 700 10 650 30 600 4 1250 850 0.5 30 680 — — 680 10650 20 450 5 C 1250 850 0.1 60 700 700 5 — — — 60 450 6 D 1250 850 0.580 680 680 5 — — — 80 450 7 E 1250 850 0.5 70 710 710 5 — — — 80 450 8 F1250 850 0.5 60 700 700 5 — — — 70 450 9 G 1250 850 0.5 80 690 690 5 — —— 80 450 10 H 1250 850 0.5 60 680 680 5 — — — 60 450 11 I 1250 850 0.160 690 — — 690 10 650 60 450 12 J 1250 850 0.1 80 700 — — 700 10 650 60450 13 K 1250 850 0.1 80 680 680 5 680 10 640 80 450 14 L 1250 850 0.360 690 690 5 690 10 650 60 450 15 H 1250 750 0.5 50 620 620 5 620 10 58060 450 16 1250 850 3.0 50 680 680 12  — — — 70 450 17 1250 850 0.5 30680 680 5 680 10 650 60 450 18 1250 850 0.5 60 600 600 5 — — — 70 450 191250 850 0.5 60 700 — — — — — 70 450

For the resulting hot-rolled steel strip (hot-rolled steel sheet), themicrostructure, tensile properties, strain age hardenability, and holeexpanding property were determined. Press formability was evaluated interms of elongation El (ductility), TS×El balance and hole expandingratio λ. Test methods were as follows.

(1) Microstructure

A test piece was sampled from each of the resultant hot-rolled sheets,and the microstructure of the cross-section (section C) perpendicular tothe rolling direction of the steel sheet was observed with an opticalmicroscope and a scanning electron microscope. The volume ratios of theferrite phase, the bainite phase, and the martensite phase in the steelsheet were determined with an image analyzer using a photograph of thecross-sectional structure at a magnification of 1,000. The volume ratiosof the retained austenite phase were determined by polishing the steelsheet to the central plane in the thickness direction, and by measuringdiffraction X-ray intensities at the central plane. Mo Kα-rays were usedas incident X-rays, the ratios of the diffraction X-ray strengths of theplanes {200}, {220} and {311} of the retained austenite phase to thediffraction X-ray strengths of the planes {110}, {200} and {211} of theferrite phase, respectively, were determined, and the volume ratio ofthe retained austenite was determined from the average of these ratios.

(2) Tensile Properties

JIS No. 5 tensile test pieces were sampled from the resultant hot-rolledsheets, and a tensile test was carried out in accordance with JIS Z 2241to determine the yield strength YS, the tensile strength TS, and theelongation El.

(3) Strain Age Hardenability

JIS No. 5 test pieces were sampled in the rolling direction from theresultant hot-rolled steel sheets. A plastic deformation of 5% wasapplied as a pre-deformation (tensile prestrain). After a heat treatmentat 250° C. for 20 minutes, a tensile test was carried out to determinetensile properties (yield stress YS_(TH) and tensile strength TS_(HT))and to calculate ΔYS=YS_(TH)−YS, and ΔTS=TS_(HT)−TS, wherein YS_(TH) andTS_(HT) were yield stress and tensile strength after thepre-deformation/heat treatment, and YS and TS were yield stress andtensile strength of the hot-rolled steel sheets.

(4) Hole Expanding Property

A hole was formed by punching a test piece sampled from the resultanthot-rolled sheet in accordance with Japan Iron and Steel FederationStandard (JFS T 1001-1996) with a punch having a diameter of 10 mm.Then, the hole was expanded with a conical punch having a vertical angleof 60° so that burrs were produced on the outside until cracks passingthrough the thickness form, thereby determining the hole expanding ratioλ. The hole expanding ratio λ was calculated by the formula: λ(%)={(d−d₀)/d₀}×100, where d₀ is initial hole diameter (punch diameter),and d is inner hole diameter upon occurrence of cracks.

The results are shown in Table 3. TABLE 3 HOLE EX- PAN- MICROSTRUCTURESION PRI- HOLE MARY PROPERTIES EX- PHASE SECONDARY PHASE AFTER PRE-PAND- F VOL- A VOL- VOL- HOT-ROLLED SHEET DEFORMATION - STRAIN AGE INGUME UME UME PROPERTIES HEAT HARDENING RA- STEEL RA- RA- OTHER RA-TENSILE PROPERTIES TREATMENT PROPERTIES TIO SHEET STEEL TIO TIO PHASESTIO YS TS El TS × El YS_(HT) TS_(HT) ΔYS ΔTS λ RE- NO. NO. % % KIND* %(MPa) (MPa) % MPa % MPa MPa MPa MPa % MARKS 1 A 75 8 B, M 25 470 620 3421080 715 790 245 170 140 EXAM- PLE 2 B 80 11 B, M 20 490 650 33 21450750 830 260 180 135 EXAM- PLE 3 75 — P 25 660 720 15 10800 730 760 70 4070 COMP. EX. 4 76 — P, B 24 600 660 16 10560 660 695 60 35 60 COMP. EX.5 C 78 9 B, M 22 490 650 33 21450 730 810 240 160 145 EXAM- PLE 6 D 75 9B, M 25 500 660 32 21120 745 825 245 165 140 EXAM- PLE 7 E 80 8 B, M 20410 540 39 21060 715 550 305 10 60 COMP. EX. 8 F 81 10 B, M 19 470 62034 21080 675 750 205 130 140 EXAM- PLE 9 G 80 9 B, M 20 460 610 35 21350690 765 230 155 135 EXAM- PLE 10 H 80 9 B, M 20 490 650 33 21450 750 830260 180 135 EXAM- PLE 11 I 81 8 B, M 19 470 620 34 21080 675 750 205 130140 EXAM- PLE 12 J 78 10 B, M 22 500 660 32 21120 745 825 245 165 140EXAM- PLE 13 K 80 8 B, M 20 470 620 34 21080 715 790 245 170 140 EXAM-PLE 14 L 75 10 B, M 25 500 660 32 21120 745 825 245 165 140 EXAM- PLE 15H 80 — P, B 20 600 660 16 10560 660 695 60 35 60 COMP. EX. 16 80 — P 20590 650 15 9750 660 690 70 40 70 COMP. EX. 17 80 — P, B 20 610 670 149380 670 705 60 35 70 COMP. EX. 18 80 — P, 20 580 640 17 10880 650 67570 35 60 COMP. EX. 19 78 — P, B 22 590 650 15 9750 650 690 60 40 70COMP. EX.*F: FERRITE,A: AUSTENITE,M: MARTENSITE,P: PEARLITE,B: BAINITE

All Examples according to the present invention show a high elongationEl, a high strength/ductility balance (TS×El), and a high hole expandingratio λ, suggesting excellent stretch flanging formability. In addition,all Examples according to the present invention show a very large ΔTS,suggesting that these samples had excellent strain age hardenability.Comparative Examples outside the scope of the present invention, incontrast, suggest that the samples have a low elongation El, a smallhole expanding ratio λ, a low ΔTS, and decreased press formability andstrain age hardenability.

Example 2

Molten steels having the compositions shown in Table 4 were made in aconverter and cast into steel slabs by a continuous casting process.Each of these steel slabs were reheated, and hot-rolled under conditionsshown in Table 5 into a hot-rolled steel strip (hot-rolled sheet) havinga thickness of 2.0 mm. The hot-rolled steel strip was temper-rolled at areduction of 1.0%. TABLE 1 STEEL COMPOSITION (wt. %) NO. C Si Mn P S AlN Cu Ni Cr, Mo, Nb, Ti, V A 0.09 1.45 1.05 0.01 0.003 0.034 0.002 1.52 —— — B 0.12 1.50 1.20 0.01 0.002 0.030 0.002 1.43 0.65 Mo: 0.32 — C 0.101.48 1.35 0.01 0.002 0.028 0.002 1.25 0.52 Cr: 0.53 — D 0.15 1.53 1.450.01 0.003 0.033 0.002 1.33 0.44 — Nb: 0.01, Ti: 0.01, V: 0.01 E 0.121.48 1.55 0.01 0.005 0.032 0.002 0.15 — — — F 0.11 1.50 1.08 0.01 0.0040.032 0.002 0.68 — — — G 0.13 1.52 1.22 0.01 0.004 0.030 0.002 0.98 — —— H 0.12 1.42 1.22 0.01 0.003 0.033 0.002 1.55 0.62 — — I 0.11 1.52 1.520.01 0.003 0.031 0.002 1.49 — Cr: 0.15, — Mo: 0.12 J 0.13 1.43 1.48 0.010.003 0.028 0.002 1.43 — Mo: 0.21 — K 0.15 1.58 1.05 0.01 0.003 0.0300.002 1.52 — — Nb: 0.01 L 0.14 1.60 1.21 0.01 0.003 0.028 0.002 1.48 —Cr: 0.11 Ti: 0.01

TABLE 1 STEEL COMPOSITION (wt. %) NO. C Si Mn P S Al N Cu Ni Cr, Mo, Nb,Ti, V A 0.09 1.45 1.05 0.01 0.003 0.034 0.002 1.52 — — — B 0.12 1.501.20 0.01 0.002 0.030 0.002 1.43 0.65 Mo: 0.32 — C 0.10 1.48 1.35 0.010.002 0.028 0.002 1.25 0.52 Cr: 0.53 — D 0.15 1.53 1.45 0.01 0.003 0.0330.002 1.33 0.44 — Nb: 0.01, Ti: 0.01, V: 0.01 E 0.12 1.48 1.55 0.010.005 0.032 0.002 0.15 — — — F 0.11 1.50 1.08 0.01 0.004 0.032 0.0020.68 — — — G 0.13 1.52 1.22 0.01 0.004 0.030 0.002 0.98 — — — H 0.121.42 1.22 0.01 0.003 0.033 0.002 1.55 0.62 — — I 0.11 1.52 1.52 0.010.003 0.031 0.002 1.49 — Cr: 0.15, — Mo: 0.12 J 0.13 1.43 1.48 0.010.003 0.028 0.002 1.43 — Mo: 0.21 — K 0.15 1.58 1.05 0.01 0.003 0.0300.002 1.52 — — Nb: 0.01 L 0.14 1.60 1.21 0.01 0.003 0.028 0.002 1.48 —Cr: 0.11 Ti: 0.01

For the resultant hot-rolled steel strip (hot-rolled steel sheet), themicrostructure, the tensile properties, the strain age hardenability,and the hole expanding ratio were determined as in Example 1. Pressformability was evaluated in terms of elongation El (ductility), TS×Elbalance and the hole expanding ratio λ.

The results obtained are shown in Table 6. TABLE 6 HOLE EX- PAN-MICROSTRUCTURE SION PRI- HOLE MARY PROPERTIES EX- PHASE SECONDARY PHASEAFTER PRE- PAND- F VOL- A VOL- VOL- HOT-ROLLED SHEET DEFORMATION -STRAIN AGE ING UME UME UME PROPERTIES HEAT HARDENING RA- STEEL RA- RA-OTHER RA- TENSILE PROPERTIES TREATMENT PROPERTIES TIO SHEET STEEL TIOTIO PHASES TIO YS TS El TS × El YS_(HT) TS_(HT) ΔYS ΔTS λ RE- NO. NO. %% KIND* % (MPa) (MPa) % MPa % MPa MPa MPa MPa % MARKS 2-1  2A 76 8 B, M24 460 610 35 21350 695 760 235 150 135 EXAM- PLE 2-2  2B 79 9 B, M 21480 640 33 21120 730 800 250 160 140 EXAM- PLE 2-3  76 — P 24 650 710 1510650 700 730 50 20 70 COMP. EX. 2-4  75 — P, B 25 590 650 14  9100 635665 45 15 65 COMP. EX. 2-5  2C 76 9 B, M 24 480 630 34 21420 715 785 235155 140 EXAM- PLE 2-6  2D 78 8 B, M 22 490 650 33 21450 725 810 235 160135 EXAM- PLE 2-7  2E 80 7 B, M 20 390 510 42 21420 620 670 230 160 130EXAM- PLE 2-8  2F 81 9 B, M 19 450 590 36 21240 660 730 210 140 135EXAM- PLE 2-9  2G 79 10 B, M 21 450 600 36 21600 570 630 120 30 65 COMP.EX. 2-10 2H 78 10 B, M 22 480 630 34 21420 715 785 235 155 130 EXAM- PLE2-11 2I 80 8 B, M 20 460 610 35 21350 695 760 235 150 135 EXAM- PLE 2-122J 79 9 B, M 21 450 590 36 21240 660 730 210 140 130 EXAM- PLE 2-13 2K80 9 B, M 20 460 600 35 21000 670 750 200 150 140 EXAM- PLE 2-14 2L 81 8B, H 19 470 620 34 21080 670 780 200 160 135 EXAM- PLE*F: FERRITE,A: AUSTENITE,M: MARTENSITE,P: PEARLITE,B: BAINITE

All Examples according to the present invention showed a high elongationEl, a high strength-ductility balance (TS×El) having excellent pressformatility, and further showed a very large ΔTS, suggesting that thesesamples had excellent strain age hardenability. Comparative Examplesoutside the scope of the present invention, in contrast, suggest thatthe samples had a low elongation El, a low ΔTS, and decreased pressformability and strain age hardenability.

Example 3

Molten steels having the composition shown in Table 7 were made in aconverter and cast into steel slabs by a continuous casting process.Then, each of these steel slabs was reheated to 1,250° C., andhot-rolled in a hot rolling step of hot rolling at a finish rolling endtemperature of 900° C. and a coiling temperature of 600° C. into ahot-rolled steel strip (hot-rolled sheet) having a thickness of 4.0 mm.Then, the hot-rolled steel strip (hot-rolled sheet) was subjected to acold rolling step of pickling and cold-rolling into cold rolled steelstrip (cold-rolled sheet) having a thickness of 1.2 mm. Thereafter, thecold-rolled steel strip (cold-rolled sheet) was subjected torecrystallization annealing step comprising heating and soakingtreatment and a subsequent retaining treatment under the conditionsshown in Table 8 on the continuous annealing line to obtain cold-rolledannealed sheet. The resultant steel strip (cold-rolled annealed sheet)was further temper-rolled at an reduction of 0.8%. TABLE 7TRANSFORMATION STEEL COMPOSITION (wt. %) POINT (° C.) NO. C Si Mn P S AlN Cu Ni Cr Mo Nb Ti V Ac1 Ac3 3A 0.10 1.20 1.42 0.01 0.003 0.032 0.0021.51 — — — — — — 725 875 3B 0.11 1.10 1.51 0.01 0.002 0.033 0.002 1.450.63 — 0.11 — — — 715 875 3C 0.11 1.32 1.33 0.01 0.004 0.025 0.002 1.200.52 0.12 — — — — 725 880 3D 0.10 1.06 1.48 0.01 0.003 0.022 0.002 1.390.43 — — 0.01 0.01 0.01 720 870 3E 0.09 1.25 1.36 0.01 0.004 0.029 0.0020.22 — — — — — — 730 860 3F 0.10 1.08 1.45 0.01 0.001 0.030 0.002 0.75 —— — — — — 720 880 3G 0.11 1.15 1.52 0.01 0.002 0.033 0.002 0.96 — — — —— — 725 875 3H 0.10 1.10 1.55 0.01 0.002 0.025 0.002 1.22 0.66 — — — — —730 875 3I 0.11 1.09 1.48 0.01 0.001 0.033 0.002 1.36 — — 0.10 — — — 725860 3J 0.11 1.12 1.62 0.01 0.002 0.029 0.001 1.42 — 0.10 — — — — 730 8803K 0.10 1.25 1.39 0.01 0.002 0.032 0.002 1.38 — — — 0.01 — — 720 870 3L0.09 1.10 1.45 0.01 0.003 0.025 0.002 1.29 — — — — 0.01 — 725 865 3M0.10 1.35 1.50 0.01 0.002 0.030 0.002 1.44 — — — — — 0.01 730 875 3N0.11 1.26 1.46 0.01 0.001 0.028 0.001 1.33 0.52 0.12 0.11 0.01 0.01 0.01725 865

TABLE 7 TRANSFORMATION STEEL COMPOSITION (wt. %) POINT (° C.) NO. C SiMn P S Al N Cu Ni Cr Mo Nb Ti V Ac1 Ac3 3A 0.10 1.20 1.42 0.01 0.0030.032 0.002 1.51 — — — — — — 725 875 3B 0.11 1.10 1.51 0.01 0.002 0.0330.002 1.45 0.63 — 0.11 — — — 715 875 3C 0.11 1.32 1.33 0.01 0.004 0.0250.002 1.20 0.52 0.12 — — — — 725 880 3D 0.10 1.06 1.48 0.01 0.003 0.0220.002 1.39 0.43 — — 0.01 0.01 0.01 720 870 3E 0.09 1.25 1.36 0.01 0.0040.029 0.002 0.22 — — — — — — 730 860 3F 0.10 1.08 1.45 0.01 0.001 0.0300.002 0.75 — — — — — — 720 880 3G 0.11 1.15 1.52 0.01 0.002 0.033 0.0020.96 — — — — — — 725 875 3H 0.10 1.10 1.55 0.01 0.002 0.025 0.002 1.220.66 — — — — — 730 875 3I 0.11 1.09 1.48 0.01 0.001 0.033 0.002 1.36 — —0.10 — — — 725 860 3J 0.11 1.12 1.62 0.01 0.002 0.029 0.001 1.42 — 0.10— — — — 730 880 3K 0.10 1.25 1.39 0.01 0.002 0.032 0.002 1.38 — — — 0.01— — 720 870 3L 0.09 1.10 1.45 0.01 0.003 0.025 0.002 1.29 — — — — 0.01 —725 865 3M 0.10 1.35 1.50 0.01 0.002 0.030 0.002 1.44 — — — — — 0.01 730875 3N 0.11 1.26 1.46 0.01 0.001 0.028 0.001 1.33 0.52 0.12 0.11 0.010.01 0.01 725 865

A test piece was sampled from the resultant steel strip, and themicrostructure, tensile properties, the strain age hardenability, andthe hole expanding property were investigated, as in Example 1. Thepress formability was evaluated in terms of the elongation El(ductility), strength-elongation balance TS×El, and the hole expandingratio, as in Example 1.

(1) Microstructure

A test piece was sampled from each of the resultant steel sheets, andthe microstructure of the cross-section (section L) in the rollingdirection of the steel sheet was observed with an optical microscope anda scanning electron microscope. The volume ratios of the ferrite,bainite, and martensite phases in the steel sheet were determined, as inExample 1, by image analysis using a photograph of the cross-sectionalstructure at a magnification of 1,000. The amount of the retainedaustenite was determined, as in Example 1, by polishing the steel sheetto the central plane in the thickness direction and by measuringdiffraction X-ray intensities at the central plane. The incident X-ray,the planes of the ferrite phase, and the planes of retained austeniteused were the same as those in Example 1.

(2) Tensile Properties

JIS No. 5 tensile test pieces were sampled from the resultant steelstrips in the direction perpendicular to the rolling direction, and atensile test was carried out, as in Example 1, in accordance with JIS Z2241 to determine yield strength YS, tensile strength TS, and elongationEl.

(3) Strain Age Hardenability

JIS No. 5 test pieces were sampled in the direction perpendicular to therolling direction from the resultant steel strips (cold-rolled annealedsheets). A plastic deformation of 5% was applied as a pre-deformation(tensile prestrain), as in Example 1. After a heat treatment at 250° C.for 20 minutes, a tensile test was carried out to determine tensileproperties (yield stress YS_(HT), and tensile strength TS_(HT)) and tocalculate ΔYS=YS_(HT)−YS, and ΔTS=TS_(HT)−TS, wherein YS_(HT) andTS_(HT) were yield stress and tensile strength after thepre-deformation—heat treatment, and YS and TS were yield stress andtensile strength of the steel strips (cold-rolled annealed sheets).

(4) Hole Expanding Property

A hole was formed by punching a test piece sampled from the resultantsteel strip in accordance with Japan Iron and Steel Federation StandardJFS T 1001-1996 with a punch having a diameter of 10 mm. Then, the holewas expanded with a conical punch having a vertical angle of 60° so thatburrs were produced on the outside until cracks passing through thethickness form, thereby determining the hole expanding ratio λ, as inExample 1.

The results are shown in Table 9. TABLE 9 MICROSTRUCTURE SECONDARY PHASECOLD-ROLLED FERRITE RETAINED SHEET PROPERTIES STEEL VOLUME AUSTENITEVOLUME TENSILE PROPERTIES SHEET STEEL RATIO VOLUME RATIO YS TS NO. NO.(%) KIND RATIO % (%) (MPa) (MPa) EI (%) TS × EI 3-1  3A 90 A, B 6 10 475630 34 21420 3-2  3B 92 A, B 5 8 500 660 32 21120 3-3  0 P, B, M 0 100690 730 11 8030 3-4  100 — 0 0 650 670 11 7370 3-5  3C 92 A, B 5 8 490650 33 21450 3-6  3D 91 A, B 5 9 500 670 32 21440 3-7  3E 93 A, B 3 7400 530 40 21200 3-8  3F 94 A, B 4 6 450 590 36 21240 3-9  3G 93 A, B 57 460 610 35 21350 3-10 3H 90 A, B 6 10 465 620 34 21080 3-11 3I 92 A, B5 8 460 610 34 20740 3-12 3J 90 A, B 6 10 500 660 32 21120 3-13 3K 92 A,B 6 8 480 640 33 21120 3-14 3L 91 A, B 5 9 470 630 33 20790 3-15 3M 90A, B 5 10 475 630 34 21420 3-15 3N 92 A, B 4 8 460 610 34 20740 3-17 3A90 P 0 10 510 600 28 16800 3-18 91 B 0 9 540 630 25 15750 3-19 90 M 0 10420 650 27 17550 3-20 92 M 0 8 430 640 28 17920 HOLE PROPERTIES AFTERSTRAIN AGE EXPANSION STEEL PRE-DEFORMATION - HARDENING HOLE SHEET STEELHEAT TREATMENT PROPERTIES EXPANDING NO. NO. YS_(HT) (MPa) TS_(HT) (MPa)ΔYS (MPa) ΔTS (MPa) RATIO λ % REMARKS 3-1  3A 710 790 235 160 140EXAMPLE 3-2  3B 750 830 250 170 135 EXAMPLE 3-3  740 760 50 30 60 COMP.EX. 3-4  690 695 40 25 130 COMP. EX. 3-5  3C 730 810 240 160 135 EXAMPLE3-6  3D 750 825 250 155 130 EXAMPLE 3-7  3E 500 550 100 20 50 COMP. EX.3-8  3F 670 740 220 150 145 EXAMPLE 3-9  3G 690 765 230 155 140 EXAMPLE3-10 3H 700 780 235 160 130 EXAMPLE 3-11 3I 705 780 245 170 135 EXAMPLE3-12 3J 740 820 240 160 130 EXAMPLE 3-13 3K 730 810 250 170 130 EXAMPLE3-14 3L 720 795 250 165 135 EXAMPLE 3-15 3M 715 790 240 160 140 EXAMPLE3-15 3N 705 780 245 170 130 EXAMPLE 3-17 3A 590 650 80 50 70 COMP. EX.3-18 605 670 65 40 120 COMP. EX. 3-19 725 805 305 155 125 COMP. EX. 3-20720 800 290 160 120 COMP. EX.F: FERRITE,A: RETAINED AUSTENITE,M: MARTENSITE,P: PEARLITE,B: BAINITE

All Examples according to the present invention are cold-rolled steelsheets having a high elongation El, a high strength-elongation balanceTS×El, a high hole expanding ratio λ, and excellent press formabilityincluding stretch flanging formability. In addition, Examples accordingto the present invention each show a very large ΔTS, suggesting that thesamples have excellent strain age hardenability. Comparative Examplesoutside the scope of the present invention, in contrast, suggest thatthe samples each have a low elongation El, a low TS−El, a small holeexpanding ratio λ, a low ΔTS, and decreased press formability and strainage hardenability.

Example 4

Molten steels having the compositions shown in Table 10 were made in aconverter and cast into steel slabs by a continuous casting process.Each of these steel slabs were reheated to 1,250° C., and hot-rolled bya hot rolling step of hot rolling with a finish rolling end temperatureof 900° C. and a coiling temperature of 600° C. into a hot-rolled steelstrip (hot-rolled sheet) having a thickness of 4.0 mm. Then, thehot-rolled steel strip (hot-rolled sheet) was subjected to a coldrolling step of pickling and cold-rolling into a cold rolled steel strip(cold-rolled sheet) having a thickness of 1.2 mm. Thereafter, thecold-rolled steel strip (cold-rolled sheet) was subjected torecrystallization annealing step comprising a heating and soakingtreatment and a subsequent retaining treatment under the conditionsshown in Table 11 on a continuous annealing line to obtain cold-rolledannealed sheet. The resultant steel strip (cold-rolled annealed sheet)was further temper-rolled at an reduction of 0.8%. TABLE 10TRANSFORMATION STEEL COMPOSITION (wt. %) POINT (° C.) NO. C Si Mn P S AlN Mo Cr W Nb Ti V Ac1 Ac3 4A 0.10 1.21 1.45 0.01 0.003 0.032 0.002 0.450.15 — — — — 740 880 4B 0.11 1.12 1.52 0.01 0.002 0.032 0.002 0.32 — —0.04 — 0.05 735 875 4C 0.11 1.30 1.35 0.01 0.003 0.028 0.002 0.48 — —0.05 0.03 — 740 885 4D 0.10 1.05 1.50 0.01 0.004 0.033 0.002 — — 0.54 —— — 735 875 4E 0.09 1.26 1.38 0.01 0.004 0.032 0.002 0.35 — — — 0.05 —735 880 4F 0.10 1.10 1.48 0.01 0.003 0.031 0.002 — 0.50 — 0.05 — — 730885 4G 0.11 1.16 1.53 0.01 0.004 0.032 0.002 — — — — — — 725 830 4H 0.121.20 1.52 0.01 0.002 0.028 0.002 0.35 — — — — — 740 870 4I 0.10 1.181.45 0.01 0.002 0.030 0.002 — 0.25 — — — — 735 860 4J 0.11 1.10 1.360.01 0.003 0.031 0.002 0.45 — — — — 730 860 4K 0.12 1.15 1.45 0.01 0.0010.025 0.002 0.30 — — 0.03 0.01 0.01 735 850 4L 0.11 1.08 1.50 0.01 0.0030.032 0.002 0.25 0.15 0.10 — — — 740 865

TABLE 11 RECRYSTALLIZATION ANNEALING HOT ROLLING STEP STEP FINISHHEATING ROLLING COLD ROLLING SOAKING SLAB END COILING STEP TREATMENTRETAINING STEEL REHEATING TEMP. TEMP. COLD ROLLING HEATING TREATMENTSHEET STEEL TEMP. FDT CT REDUCTION SOAKING TEMP RETENTION NO. NO. (° C.)° C. ° C. % TEMP. (° C.) (° C.) TIME (s) 4-1  4A 1250 900 600 70 800 400300 4-2  4B 1250 900 600 70 800 400 300 4-3  1250 900 600 70 980 — —4-4  1250 900 600 70 680 400 300 4-5  4C 1250 900 600 70 800 400 3004-6  4D 1250 900 600 70 800 400 300 4-7  4E 1250 900 600 70 800 400 3004-8  4F 1250 900 600 70 800 400 300 4-9  4G 1250 900 600 70 800 400 3004-10 4H 1250 900 600 70 800 400 300 4-11 4I 1250 900 600 70 800 400 3004-12 4J 1250 900 600 70 800 400 300 4-13 4K 1250 900 600 70 800 400 3004-14 4L 1250 900 600 70 800 400 300 4-15 4A 1250 900 600 70 800 250 3004-16 1250 900 600 70 800 550 300

A test piece was sampled from the resultant steel strip, and themicrostructure, the tensile properties, the strain age hardenability,and the hole expanding property were investigated, as in Example 3.

The results are shown in Table 12. TABLE 12 MICROSTRUCTURE SECONDARYPHASE COLD-ROLLED FERRITE RETAINED SHEET PROPERTIES STEEL VOLUMEAUSTENITE VOLUME TENSILE PROPERTIES SHEET STEEL RATIO VOLUME RATIO YS TSNO. NO. (%) KIND RATIO % (%) (MPa) (MPa) El (%) TS × El 4-1  4A 91 A, B6 9 470 630 34 21420 4-2  4B 92 A, B 5 8 500 660 32 21120 4-3  0 P, B, M0 100 560 740 12 8880 4-4  100 — 0 0 500 660 11 7260 4-5  4C 92 A, B 5 8480 640 33 21120 4-6  4D 94 A, B 4 6 470 630 34 21420 4-7  4E 92 A, B 58 490 650 33 21450 4-8  4F 93 A, B 4 7 470 620 34 21080 4-9  4G 94 A, B3 6 460 620 34 21080 4-10 4H 92 A, B 5 8 475 630 33 20790 4-11 4I 90 A,B 4 10 480 640 33 21120 4-12 4J 91 A, B 5 9 485 650 32 20800 4-13 4K 92A, B 4 8 470 630 34 21420 4-14 4L 90 A, B 5 10 465 620 34 21080 4-15 4A93 M 0 7 380 630 28 17640 4-16 92 P 0 8 550 650 24 15600 HOLE PROPERTIESAFTER STRAIN AGE EXPANSION STEEL PRE-DEFORMATION - HARDENING HOLE SHEETSTEEL HEAT TREATMENT PROPERTIES EXPANDING NO. NO. YS_(HT) (MPa) TS_(HT)(MPa) ΔYS (MPa) ΔTS (MPa) RATIO λ % REMARKS 4-1  4A 700 780 230 150 140EXAMPLE 4-2  4B 740 820 240 160 130 EXAMPLE 4-3  680 760 120 20 60 COMP.EX. 4-4  610 675 110 15 130 COMP. EX. 4-5  4C 710 790 230 150 120EXAMPLE 4-6  4D 700 775 230 145 130 EXAMPLE 4-7  4E 720 800 230 150 120EXAMPLE 4-8  4F 680 760 210 140 120 EXAMPLE 4-9  4G 570 630 110 10 60COMP. EX. 4-10 4H 710 790 235 160 130 EXAMPLE 4-11 4I 725 805 245 165120 EXAMPLE 4-12 4J 730 810 245 160 120 EXAMPLE 4-13 4K 710 790 240 160130 EXAMPLE 4-14 4L 700 775 235 155 120 EXAMPLE 4-15 4A 710 790 330 160110 COMP. EX. 4-16 620 680 70 30 70 COMP. EX.F: FERRITE,A: RETAINED AUSTENITE,M: MARTENSITE,P: PEARLITE,B: BAINITE

All Examples according to the present invention show a high elongationEl, a high strength-ductility balance TS×El, and a high hole expandingratio λ, suggesting that the samples have excellent press formabilityincluding stretch flanging formability. In addition, Examples accordingto the present invention show a very large ΔTS, suggesting that thesamples have excellent strain age hardenability. Comparative Examplesoutside the scope of the present invention, in contrast, suggest thatthe samples have a low elongation El, a low TS×El, a small holeexpanding ratio λ, a low ΔTS, and decreased press formability and strainage hardenability.

Example 5

Molten steels having the compositions shown in Table 13 were made in aconverter and cast into steel slabs by a continuous casting process.These slabs were hot-rolled under the conditions shown in Table 14 intohot-rolled steel strips (hot-rolled sheets).

After pickling, each of these hot-rolled steel strips (hot-rolledsheets) was subjected to a primary heat treatment step on a continuousannealing line (CAL) under the conditions shown in Table 14 and asecondary heat treatment step on a continuous hot-dip galvanizing line(CGL) under the conditions shown in Table 14. Then, the sheet wassubjected to a hot-dip galvanizing treatment step of performing ahot-dip galvanizing which forms a hot-dip galvanizing layer on thesurfaces of the steel sheet. Then, an alloying treatment step ofalloying the hot-dip galvanizing layer was applied under the conditionsshown in Table 14. Some of the steel sheets were left as hot-dipgalvanized.

After further pickling, the hot-rolled steel strip (hot-rolled sheet)obtained by the above-mentioned hot rolling was subjected to a coldrolling step under the conditions shown in Table 14 into a cold-rolledsteel strip (cold-rolled sheet). Then, the cold-rolled steel strip(cold-rolled sheet) was subjected to a primary heat treatment step on acontinuous annealing line (CAL) under the conditions shown in Table 14.After a secondary heat treatment step on the continuous hot-dipgalvanizing line (CGL) under the conditions shown in Table 14, a hot-dipgalvanizing treatment step was performed. Then, an alloying treatmentstep was performed under the conditions shown in Table 14. Some of thesteel sheets were left as hot-dip galvanized.

Prior to the secondary heat treatment step on the continuous hot-dipgalvanizing line (CGL), some of the steel sheets after the primary heattreatment step were subjected to a pickling treatment shown in Table 14.The pickling treatment was carried out in a pickling bath on the entryside of the CGL.

The galvanizing bath temperature was within the range of 460 to 480° C.,and the temperature of the steel sheet to be dipped was within the rangeof the galvanizing bath temperature to (bath temperature+10° C.) In thealloying treatment, the sheet was reheated within the temperature rangeof 480 to 540° C., and held at the temperature for 15 to 28 seconds. Thecooling rate after the alloying treatment was 10° C./second. The platedsteel sheet was further temper rolled at a reduction of 1.0%. TABLE 13TRANS- FORMATION STEEL COMPOSITION (wt. %) POINT (° C.) NO. C Si Mn P SAl N Cu Ni Cr, Mo Nb, Ti, V Ac1 Ac3 5A 0.08 0.72 2.05 0.01 0.003 0.0320.002 1.48 — — — 715 875 5B 0.07 0.52 2.22 0.01 0.001 0.033 0.002 1.440.62 Mo: 0.15 — 720 870 5C 0.09 0.77 1.85 0.01 0.004 0.028 0.002 1.280.55 Cr: 0.15 — 725 875 5D 0.08 0.65 1.95 0.01 0.005 0.032 0.002 1.330.42 — Nb: 0.01, 715 870 Ti: 0.01, V: 0.01 5E 0.07 0.55 2.05 0.01 0.0040.033 0.002 0.14 — — — 715 875 5F 0.08 0.70 2.22 0.01 0.003 0.033 0.0020.72 — — — 715 870 5G 0.07 0.68 1.85 0.01 0.005 0.036 0.002 0.95 — — —715 875 5H 0.08 0.77 2.05 0.01 0.003 0.032 0.002 1.45 0.75 — — 715 8705I 0.09 0.80 1.85 0.01 0.002 0.028 0.002 1.29 — Cr: 0.12 — 720 875 5J0.07 0.75 2.05 0.01 0.005 0.030 0.002 1.38 — Mo: 0.15 — 715 870 5K 0.080.68 1.95 0.01 0.003 0.025 0.002 1.40 — — Nb: 0.01 720 875 5L 0.07 0.702.10 0.01 0.004 0.030 0.002 1.35 — — Ti: 0.01 715 870 5M 0.08 0.75 1.800.01 0.002 0.031 0.002 1.25 — — V: 0.01 725 870 5N 0.09 0.68 2.00 0.010.003 0.035 0.002 1.35 0.60 Cr: 0.13, Nb: 0.01, 710 875 Mo: 0.15 V: 0.01

TABLE 14 HOT ROLLING STEP FINISH ROLLING COLD ROLLING STEP PRIMARY HEATTREATMENT SLAB END COILING FINAL COLD FINAL STEP STEEL REHEATING TEMP.TEMP. THICK- ROLLING THICK- HEATING COOLING SHEET STEEL TEMP. FDT CTNESS REDUCTION NESS TEMP. RATE NO. NO. (° C.) ° C. ° C. mm % mm LINE °C. ° C./s 5-1 5A 1250 850 600 1.2 — — CAL 880 20 5-2 5B 1250 850 600 1.2— — CAL 880 20 5-3 5-4 5-5 5-6 5C 1250 850 600 1.2 — — CAL 880 20 5-7 5D1250 850 600 1.2 — — CAL 880 20 5-8 5E 1250 850 600 1.2 — — CAL 880 205-9 5F 1250 850 600 1.2 — — CAL 880 20 5-10 5G 1250 850 600 1.2 — — CAL880 20 5-11 5A 1250 850 600 4.0 70 1.2 CAL 880 20 5-12 5B 1250 850 6004.0 70 1.2 CAL 880 20 5-13 CAL 880 20 5-14 CAL 880 20 5-15 CAL 880 205-16 5C 1250 850 600 4.0 70 1.2 CAL 880 20 5-17 5D 1250 850 600 4.0 701.2 CAL 880 20 5-18 5E 1250 850 600 4.0 70 1.2 CAL 880 20 5-19 5F 1250850 600 4.0 70 1.2 CAL 880 20 5-20 5G 1250 850 600 4.0 70 1.2 CAL 880 205-21 5H 1250 850 600 4.0 70 1.2 CAL 880 20 5-22 5I 1250 850 600 4.0 701.2 CAL 880 20 5-23 5J 1250 850 600 4.0 70 1.2 CAL 880 20 5-24 5K 1250850 600 4.0 70 1.2 CAL 880 20 5-25 5L 1250 850 600 4.0 70 1.2 CAL 880 205-26 5M 1250 850 600 4.0 70 1.2 CAL 880 20 5-27 5N 1250 850 600 4.0 701.2 CAL 880 20 HOT-DIP SECONDARY HEAT GALVANIZING TREATMENT STEP COOLINGPICK- HEAT- COOL- RATE AFTER ALLOYING TEMPER STEEL LING KIND ING INGKIND GALVA- TREATMENT STEP ROLLING SHEET STEEL TREAT- OF TEMP. RATE* OFNIZING** TEMP REDUCTION NO. NO. MENT LINE ° C. ° C./s LINE ° C./s ° C. %5-1 5A YES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-2 5B — CGL 800 20 CGL 10ALLOYING 500 1.0 5-3 YES CGL 780 20 CGL 10 ALLOYING 500 1.0 5-4 CGL 98020 CGL 10 ALLOYING 500 1.0 5-5 CGL 650 20 CGL 10 ALLOYING 500 1.0 5-6 5CYES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-7 5D YES CGL 820 20 CGL 10ALLOYING 500 1.0 5-8 5E YES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-9 5FYES CGL 780 20 CGL 10 NON- — 1.0 ALLOYING 5-10 5G YES CGL 800 20 CGL 10ALLOYING 500 1.0 5-11 5A YES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-12 5B— CGL 820 20 CGL 10 ALLOYING 500 1.0 5-13 YES CGL 800 20 CGL 10 ALLOYING500 1.0 5-14 YES CGL 980 20 CGL 10 ALLOYING 500 1.0 5-15 YES CGL 680 20CGL 10 ALLOYING 500 1.0 5-16 5C YES CGL 800 20 CGL 10 ALLOYING 500 1.05-17 5D YES CGL 800 20 CGL 10 NON- — 1.0 ALLOYING 5-18 5E YES CGL 780 20CGL 10 ALLOYING 500 1.0 5-19 5F YES CGL 800 20 CGL 10 ALLOYING 500 1.05-20 5G YES CGL 820 20 CGL 10 ALLOYING 500 1.0 5-21 5H YES CGL 800 20CGL 10 ALLOYING 500 1.0 5-22 5I YES CGL 800 20 CGL 10 ALLOYING 500 1.05-23 5J YES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-24 5K YES CGL 800 20CGL 10 ALLOYING 500 1.0 5-25 5L YES CGL 800 20 CGL 10 ALLOYING 500 1.05-26 5M YES CGL 800 20 CGL 10 ALLOYING 500 1.0 5-27 5N YES CGL 800 20CGL 10 ALLOYING 500 1.0*COOLING RATE UNTIL 480° C.**COOLING RATE UNTIL 300° C.

For the hot-dip galvanized steel sheet (steel strip) obtained throughthe above-mentioned steps, the microstructure, the tensile properties,the strain age hardenability, and the hole expanding ratio weredetermined, as in Example 1. Press formability was evaluated in terms ofelongation El (ductility), and hole expanding ratio.

(1) Microstructure

The microstructure of the cross-section (section L) in the rollingdirection of the steel sheet was observed with an optical microscope anda scanning electron microscope. The volume ratios of the ferrite phase,lath martensite phase, tempered martensite phase, and martensite phasewere determined, as in Example 1, by image analysis using a photographof cross-sectional structure at a magnification of 1,000. The amount ofretained austenite was determined, as in Example 1, by polishing thesteel sheet to the central plane in the thickness direction and bymeasuring diffraction X-ray intensities at the central plane. Theincident X-ray, the planes of the ferrite phase, and the planes ofretained austenite used were the same as those in Example 1.

(2) Tensile Properties

JIS No. 5 tensile test pieces were sampled from the resultant steelstrips in the direction perpendicular to the rolling direction, and atensile test was carried out in accordance with JIS Z 2241 to determinethe yield strength YS, the tensile strength TS, and the elongation El,as in Example 1.

(3) Strain Age Hardenability

JIS No. 5 test pieces were sampled from the resultant steel strips inthe direction perpendicular to the rolling direction, and a plasticdeformation of 5% was applied as a pre-deformation (tensile prestrain),as in Example 1. After a heat treatment at 250° C. for 20 minutes, atensile test was carried out to determine tensile properties (yieldstress YS_(TH), and tensile strength TS_(HT)) and to calculateΔYS=YS_(TH)−YS, and ΔTS=TS_(HT)−TS, wherein YS_(TH) and TS_(HT) wereyield stress and tensile strength after the pre-deformation—heattreatment, and YS and TS were yield stress and tensile strength of thesteel strips.

(4) Hole Expanding Ratio

A hole was formed by punching a test piece sampled from the resultantsteel strip in accordance with Japan Iron and Steel Federation StandardJFS T 1001-1996 with a punch having a diameter of 10 mm. Then, the holewas expanded with a conical punch having a vertical angle of 60° C. sothat burrs were produced on the outside until cracks passing through thethickness form, thereby determining the hole expanding ratio λ, as inExample 1.

The results are shown in Table 15. TABLE 15 MICROSTRUCTURE PRIMARY PHASESECONDARY PHASE TEMPERED RETAINED PLATED SHEET PROPERTIES STEEL FERRITEMARTENSITE AUSTENITE TENSILE PROPERTIES SHEET STEEL VOLUME VOLUME VOLUMEVOLUME VOLUME YS TS El TS × El NO. NO. RATIO % RATIO % RATIO % KIND*RATIO % RATIO % (MPa) (MPa) (%) (Mpa %) 5-1 5A 57 35 92 A, B, 5 8 470620 34 21080 5-2 5B 52 40 92 A, B, 4 8 480 640 33 21120 5-3 51 40 91 A,B, 5 9 470 620 34 21080 5-4  0  0 0 M, P, B 0 100 670 710 11 7810 5-5 6040 100 — 0 0 620 650 12 7800 5-6 5C 58 35 93 A, B 4 7 470 630 34 214205-7 5D 57 35 92 A, B 5 8 490 650 33 21450 5-8 5E 53 40 93 A, B 7 7 380510 42 21420 5-9 5F 37 55 92 A, B 4 8 430 570 37 21090 5-10 5G 53 40 93A, B 5 7 450 590 36 21240 5-11 5A 57 35 92 A, B 7 8 470 630 34 214205-12 5B 52 40 92 A, B 5 8 500 660 32 21120 5-13 53 40 93 A, B 6 7 480640 33 21120 5-14  0  0 0 M, P, B 0 100 680 720 12 8640 5-15 65 35 100 —0 0 620 660 11 7260 5-16 5C 52 40 92 A, B 4 8 490 650 33 21450 5-17 5D53 40 93 A, B 5 7 500 660 32 21120 5-18 5E 48 45 93 A, B 4 7 390 520 4121320 5-19 5F 44 50 94 A, B 5 6 440 580 37 21460 5-20 5G 57 35 92 A, B 58 450 600 35 21000 5-21 5H 51 40 91 A, B 5 9 445 590 35 20650 5-22 5I 5535 90 A, B 5 10 460 610 34 20740 5-23 5J 52 40 92 A, B 4 8 450 600 3521000 5-24 5K 53 40 93 A, B 5 7 470 620 34 21080 5-25 5L 56 35 91 A, B 69 475 630 33 20790 5-26 5M 60 30 90 A, B 5 10 460 610 34 20740 5-27 5N52 40 92 A, B 4 8 455 600 35 21000 PROPERTIES AFTER PRE- DEFORMATION -STRAIN AGE HOLE HEAT HARDENING EXPANSION STEEL TREATMENT PROPERTIES HOLESHEET STEEL YS_(HT) TS_(HT) ΔYS ΔTS EXPANDING NO. NO. (MPa) (MPa) (MPa)(MPa) RATIO λ % REMARKS 5-1 5A 700 775 230 155 140 EXAMPLE 5-2 5B 725805 245 165 135 EXAMPLE 5-3 710 785 240 165 135 EXAMPLE 5-4 710 740 4030 65 COMP. EX. 5-5 650 675 30 25 130 COMP. EX. 5-6 5C 710 785 240 155135 EXAMPLE 5-7 5D 725 805 235 155 130 EXAMPLE 5-8 5E 480 530 100 20 60COMP. EX. 5-9 5F 650 720 220 150 140 EXAMPLE 5-10 5G 675 745 225 155 135EXAMPLE 5-11 5A 715 790 245 160 145 EXAMPLE 5-12 5B 750 830 250 170 140EXAMPLE 5-13 730 810 250 170 140 EXAMPLE 5-14 720 750 40 30 70 COMP. EX.5-15 650 685 30 25 60 COMP. EX. 5-16 5C 730 810 240 160 140 EXAMPLE 5-175D 735 815 235 155 135 EXAMPLE 5-18 5E 490 540 100 20 60 COMP. EX. 5-195F 655 725 215 145 135 EXAMPLE 5-20 5G 675 750 225 150 140 EXAMPLE 5-215H 680 755 235 165 130 EXAMPLE 5-22 5I 695 770 235 160 135 EXAMPLE 5-235J 680 755 230 155 130 EXAMPLE 5-24 5K 710 780 240 160 130 EXAMPLE 5-255L 720 795 245 165 135 EXAMPLE 5-26 5M 695 770 235 160 130 EXAMPLE 5-275N 680 755 225 155 130 EXAMPLE*M: MARTENSITE,P: PEARLITE,B: BAINITE,A: RETAINED AUSTENITE

All Examples according to the present invention each show a highelongation El and a high hole expanding ratio λ, suggesting that thesamples are hot-dip galvanized steel sheets having an excellent stretchflanging formability. In addition, Examples according to the presentinvention showed a very large ΔTS, suggesting that the samples are steelsheets having excellent strain age hardenability. Comparative Examplesoutside the scope of the invention, in contrast, suggest that thesamples are steel sheets having a low elongation El, a small holeexpanding ratio λ, a low ΔTS, and decreased press formability and strainage hardenability.

Example 6

Molten steels having the compositions shown in Table 16 was made in aconverter and cast into steel slabs by a continuous casting process.Each of these steel slabs were reheated to 1,250° C., and hot-rolled bya hot rolling step of hot rolling with a finish rolling end temperatureof 900° C. and a coiling temperature of 600° C. into hot-rolled steelstrip (hot-rolled sheet) having a thickness of 4.0 mm. Then, thehot-rolled steel strip (hot-rolled sheet) was subjected to a coldrolling step of pickling and cold-rolling into cold-rolled steel strip(cold-rolled sheet) having a thickness of 1.2 mm. Then, the cold-rolledsteel strip (cold-rolled sheet) was subjected to a primary heattreatment step on a continuous annealing line (CAL) under the conditionsshown in Table 17. Then, the sheet was subjected to a secondary heattreatment step on a continuous hot-dip galvanizing line (CGL) under theconditions shown in Table 17 and then, subjected to a hot-dipgalvanizing treatment step to form a hot-dip galvanizing layer on thesurfaces of the steel sheet. In addition, an alloying treatment step wasapplied under the conditions shown in FIG. 17. The cooling rate afterthe alloying treatment was 10° C./second. Some of the steel strips(steel sheets) were left as hot-dip galvanized. TABLE 16 TRANSFORMATIONSTEEL COMPOSITION (wt. %) POINT (° C.) NO. C Si Mn P S Al N Cr, Mo, WNb, Ti, V Ac1 Ac3 6A 0.07 0.77 2.00 0.01 0.003 0.033 0.002 Cr: 0.20, —715 870 Mo: 0.43 6B 0.08 0.55 2.22 0.01 0.001 0.033 0.002 Mo: 0.33 Nb:0.04, 720 865 V: 0.05 6C 0.08 0.75 1.80 0.01 0.004 0.020 0.002 Mo: 0.48Nb: 0.05, 725 880 Ti: 0.03 6D 0.09 0.63 1.98 0.01 0.005 0.025 0.002 W:0.54 — 715 865 6E 0.07 0.65 2.02 0.01 0.003 0.033 0.002 Mo: 0.36 Ti:0.05 715 875 6F 0.08 0.70 1.90 0.01 0.005 0.035 0.002 Cr: 0.50 Nb: 0.05715 865 6G 0.07 0.58 2.08 0.01 0.004 0.032 0.002 — — 715 865 6H 0.080.75 2.22 0.01 0.004 0.022 0.002 Mo: 0.35 — 715 870 6I 0.08 0.77 1.980.01 0.003 0.032 0.002 Cr: 0.25 — 710 860 6J 0.07 0.68 2.05 0.01 0.0020.035 0.002 Mo: 0.15, — 720 865 Cr: 0.10, W: 0.11 6K 0.09 0.70 1.98 0.010.001 0.028 0.002 Mo: 0.25, V: 0.05 715 865 Cr: 0.10

TABLE 17 HOT ROLLING STEP FINISH ROLLING COLD ROLLING STEP PRIMARY HEATTREATMENT SLAB END COILING FINAL COLD FINAL STEP STEEL REHEATING TEMP.TEMP. THICK- ROLLING THICK- HEATING COOLING SHEET STEEL TEMP. FDT CTNESS REDUCTION NESS TEMP. RATE NO. NO. (° C.) ° C. ° C. mm % mm LINE °C. ° C./s 6-1 6A 1250 850 600 1.2 — — CAL 880 20 6-2 6B 1250 850 600 1.2— — CAL 880 20 6-3 6-4 6-5 6-6 6C 1250 850 600 1.2 — — CAL 880 20 6-7 6D1250 850 600 1.2 — — CAL 880 20 6-8 6E 1250 850 600 1.2 — — CAL 880 206-9 6F 1250 850 600 1.2 — — CAL 880 20 6-10 6G 1250 850 600 1.2 — — CAL880 20 6-11 6A 1250 850 600 4.0 70 1.2 CAL 880 20 6-12 6B 1250 850 6004.0 70 1.2 CAL 880 20 6-13 CAL 880 20 6-14 CAL 880 20 6-15 CAL 880 206-16 6C 1250 850 600 4.0 70 1.2 CAL 880 20 6-17 6D 1250 850 600 4.0 701.2 CAL 880 20 6-18 6E 1250 850 600 4.0 70 1.2 CAL 880 20 6-19 6F 1250850 600 4.0 70 1.2 CAL 880 20 6-20 6G 1250 850 600 4.0 70 1.2 CAL 880 206-21 6H 1250 850 600 4.0 70 1.2 CAL 880 20 6-22 6I 1250 850 600 4.0 701.2 CAL 880 20 6-23 6J 1250 850 600 4.0 70 1.2 CAL 880 20 6-24 6K 1250850 600 4.0 70 1.2 CAL 880 20 HOT-DIP SECONDARY HEAT GALVANIZINGTREATMENT STEP COOLING PICK- HEAT- COOL- RATE AFTER ALLOYING TEMPERSTEEL LING KIND ING ING KIND GALVA- TREATMENT STEP ROLLING SHEET STEELTREAT- OF TEMP. RATE* OF NIZING** TEMP REDUCTION NO. NO. MENT LINE ° C.° C./s LINE ° C./s ° C. % 6-1 6A YES CGL 780 20 CGL 10 ALLOYING 500 1.06-2 6B — CGL 800 20 CGL 10 ALLOYING 500 1.0 6-3 YES CGL 800 20 CGL 10ALLOYING 500 1.0 6-4 CGL 980 20 CGL 10 ALLOYING 500 1.0 6-5 CGL 650 20CGL 10 ALLOYING 500 1.0 6-6 6C YES CGL 780 20 CGL 10 ALLOYING 500 1.06-7 6D YES CGL 820 20 CGL 10 ALLOYING 500 1.0 6-8 6E YES CGL 800 20 CGL10 ALLOYING 500 1.0 6-9 6F YES CGL 800 20 CGL 10 NON- — 1.0 ALLOYING6-10 6G YES CGL 800 20 CGL 10 ALLOYING 500 1.0 6-11 6A YES CGL 800 20CGL 10 ALLOYING 500 1.0 6-12 6B — CGL 820 20 CGL 10 ALLOYING 500 1.06-13 YES CGL 780 20 CGL 10 ALLOYING 500 1.0 6-14 YES CGL 980 20 CGL 10ALLOYING 500 1.0 6-15 YES CGL 680 20 CGL 10 ALLOYING 500 1.0 6-16 6C YESCGL 800 20 CGL 10 ALLOYING 500 1.0 6-17 6D YES CGL 800 20 CGL 10 NON- —1.0 ALLOYING 6-18 6E YES CGL 780 20 CGL 10 ALLOYING 500 1.0 6-19 6F YESCGL 800 20 CGL 10 ALLOYING 500 1.0 6-20 6G YES CGL 820 20 CGL 10ALLOYING 500 1.0 6-21 6H YES CGL 800 20 CGL 10 ALLOYING 500 1.0 6-22 6IYES CGL 800 20 CGL 10 ALLOYING 500 1.0 6-23 6J YES CGL 800 20 CGL 10ALLOYING 500 1.0 6-24 6K YES CGL 800 20 CGL 10 ALLOYING 500 1.0*COOLING RATE UNTIL 480° C.**COOLING RATE UNTIL 300° C.

A piece was sampled from the resultant hot-dip galvanized steel strip,and the microstructure, the tensile properties, the strain agehardenability, and the bore expanding property were investigated, as inExample 5.

The results are shown in Table 18. TABLE 18 MICROSTRUCTURE PRIMARY PHASESECONDARY PHASE TEMPERED RETAINED PLATED SHEET PROPERTIES STEEL FERRITEMARTENSITE AUSTENITE TENSILE PROPERTIES SHEET STEEL VOLUME VOLUME VOLUMEVOLUME VOLUME YS TS El TS × El NO. NO. RATIO % RATIO % RATIO % KIND*RATIO % RATIO % (MPa) (MPa) (%) (Mpa %) 6-1 6A 56 35 91 A, B 6 9 460 61035 21350 6-2 6B 52 40 92 A, B 5 8 475 630 34 21420 6-3 50 40 90 A, B 610 460 610 35 21350 6-4  0  0 0 M, P, B 0 100 660 700 11 7700 6-5 60 40100 — 0 0 620 660 12 7920 6-6 6C 47 45 92 A, B 5 8 570 620 34 21080 6-76D 53 40 93 A, B 5 7 480 640 33 21120 6-8 6E 57 35 92 A, B 6 8 390 52041 21320 6-9 6F 48 45 93 A, B 5 7 420 560 38 21280 6-10 6G 53 40 93 A, B5 7 450 590 36 21240 6-11 6A 53 40 93 A, B 5 7 465 620 34 21080 6-12 6B52 40 92 A, B 5 8 490 650 33 21450 6-13 57 35 92 A, B 5 8 475 630 3421420 6-14  0  0 0 M, P, B 0 100 650 710 12 8520 6-15 60 40 100 — 0 0610 650 11 7150 6-16 6C 53 40 93 A, B 5 7 480 640 33 21120 6-17 6D 62 3092 A, B 5 8 490 650 33 21450 6-18 6E 53 40 93 A, B 4 7 390 520 41 213206-19 6F 49 45 94 A, B 4 6 450 590 36 21240 6-20 6G 42 50 92 A, B 5 8 460610 35 21350 6-21 6H 36 55 91 A, B 5 9 470 630 34 21420 6-22 6I 40 50 90A, B 4 10 465 620 34 21080 6-23 6J 50 40 90 A, B 5 10 480 640 33 211206-24 6K 51 40 91 A, B 5 9 470 620 34 21080 PROPER- TIES AFTER PRE-STRAIN DEFOR- AGE MATION - HARD- HEAT ENING HOLE TREAT- PROPER-EXPANSION STEEL MENT TIES HOLE SHEET STEEL YS_(HT) TS_(HT) ΔYS ΔTSEXPANDING NO. NO. MPa MPa MPa MPa RATIO λ % REMARKS 6-1 6A 705 780 245170 140 EXAMPLE 6-2 6B 730 810 255 180 135 EXAMPLE 6-3 715 790 255 180135 EXAMPLE 6-4 720 730 60 30 55 COMP. EX. 6-5 660 685 40 25 125 COMP.EX. 6-6 6C 715 790 145 170 135 EXAMPLE 6-7 6D 730 810 250 170 130EXAMPLE 6-8 6E 620 685 230 165 130 EXAMPLE 6-9 6F 655 725 235 165 140EXAMPLE 6-10 6G 560 620 110 30 50 COMP. EX. 6-11 6A 720 795 255 175 145EXAMPLE 6-12 6B 755 835 265 185 140 EXAMPLE 6-13 730 810 255 180 140EXAMPLE 6-14 720 740 70 30 60 COMP. EX. 6-15 650 675 40 25 50 COMP. EX.6-16 6C 730 810 250 170 135 EXAMPLE 6-17 6D 740 820 250 170 140 EXAMPLE6-18 6E 615 680 225 160 140 EXAMPLE 6-19 6F 675 750 225 160 135 EXAMPLE6-20 6G 700 775 240 165 30 COMP. EX. 6-21 6H 710 790 240 165 120 EXAMPLE6-22 6I 705 785 240 165 120 EXAMPLE 6-23 6J 720 800 240 160 130 EXAMPLE6-24 6K 700 775 230 155 120 EXAMPLE*M: MARTENSITE,P: PEARLITE,B: BAINITE,A: RETAINED AUSTENITEAll Examples according to the present invention show a high elongationEl and a high bore expanding ratio λ, suggesting that the examples arehot-dip galvanized steel sheets having excellent press formability. Inaddition, all Examples according to the present invention show a verylarge ΔTS, suggesting that the samples are steel sheets having excellentstrain age hardenability. Comparative Examples outside the scope of theinvention, in contrast, suggest that the samples are steel sheets havinga low elongation El, a low λ, a low ΔTS, and decreased press formabilityand strain age hardenability.

According to the present invention, it is possible to stably manufacturesteel sheets (hot-rolled steel sheets, cold-rolled steel sheets andhot-dip galvanized steel sheets) in which the tensile strength isremarkably increased through a heat treatment applied after pressforming while maintaining excellent press formability, givingindustrially remarkable effects. When applying a steel sheet of thepresent invention to automotive parts, there are available advantages ofeasy press forming, high and stable parts properties after completion,and sufficient contribution to the weight reduction of the automobilebody.

1-22. (cancelled)
 23. A high-ductility hot-dip galvanized steel sheetcomprising a hot-dip galvanizing layer or an alloyed hot-dip galvanizinglayer formed on the surface of a high-ductility steel sheet excellent inpress formability and in strain age hardenability as represented by aΔTS of not less than 80 mpa, comprising a composite structure containina primary phase containing a ferrite phase and a secondary phasecontaining a retained austenite phase in a volume ratio of not less than1%. 24 A high-ductility hot-dip galvanized steel sheet comprising ahot-dip galvanizing layer or an alloyed hot-dip galvanizing layer formedon the surface of a high-ductility cold-rolled steel sheet excellent inpress formability and in strain age hardenability as represented by aΔTS of not less than 80 MPa, comprising a composite structure containinga primary ferrite phase and a secondary phase containing a retainedaustenite phase in a volume ratio of not less than 1%.
 25. Ahigh-ductility steel sheet to excellent in press formability and instrain age hardenability as represented by a ΔTS of not less than 80MPa, comprising a composite structure containing a primary phasecontaining a ferrite phase and a secondary phase containing a retainedaustenite phase in a volume ratio of not less than 1%, wherein the steelsheet is a hot-dip galvanized steel sheet having a hot-dip galvanizinglayer or an alloyed hot-dip galvanizing layer formed on a surface of thesteel sheet, and the primary phase containing a ferrite phase comprisesa ferrite phase and a tempered martensite phase.
 26. A high-ductilitysteel sheet according to claim 25, wherein the steel sheet has acomposition comprising, in weight percent, C: not more than 0.20%, Si:not more than 2.0%, Mn: not more than 3.0%, P: not more than 0.1%, S:not more than 0.02%, Al: not more than 0.3%, N: not more than 0.02%, Cu:0.5 to 3.0%, and the balance Fe and incidental impurities.
 27. Ahigh-ductility steel sheet according to claim 26, the compositionfurther comprising, in weight percent, at least one of the followingGroups A to C: Group A: Ni: not more than 2.0%; Group B: at least one ofCr and Mo: not more than 2.0% in total; and Group C: at least one of Nb,Ti, and V: not more than 0.2% in total.
 28. A high-ductility steel sheetaccording to claim 25, wherein the steel sheet has a compositioncomprising, in weight percent, C: not more than 0.20%, Si: not more than2.0%, Mn: not more than 3.0%, P: not more than 0.1%, S: not more than0.02%, Al: not more than 0.3%, N: not more than 0.02%, at least oneselected from the group consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0%and W: 0.05 to 2.0% in a total amount of not more than 2.0%, and thebalance Fe and incidental impurities.
 29. A high-ductility steel sheetaccording to claim 28, the composition further comprising, in weightpercent, at least one of Nb, Ti, and V in a total amount of not morethan 2.0%.
 30. A method of manufacturing of a high-ductility hot-dipgalvanized steel sheet excellent in press formability and in strain agehardenability as typically represented by a ΔTS of not less than 80 MPa,comprising: a primary heat-treating step of heating a steel sheet to atemperature of not less than the A_(C1) transformation point and rapidlycooling the steel sheet, the steel sheet having a compositioncontaining, in weight percent, C: not more than 0.20%, Si: not more than2.0%, Mn: not more than 3.0%, P: not more than 0.1%, S: not more than0.02%, Al: not more than 0.3%, N: not more than 0.02%, and Cu: 0.5 to3.0%; a secondary heat-treating step of heating the steel sheet to atemperature in the range of the A_(C1) transformation point to theA_(C3) transformation point; and a hot-dip galvanizing step of forming ahot-dip galvanizing layer on the surface of the steel sheet.
 31. Amethod for manufacturing a high-ductility cold-rolled steel sheetaccording to claim 30, the composition further containing, in weightpercent, at least one of the following Groups A to C: Group A: Ni: notmore than 2.0%; Group B: at least one of Cr and Mo: not more than 2.0%in total; and Group C: at least one of Nb, Ti, and V: not more than 0.2%in total.
 32. A method for manufacturing a high-ductility hot-dipgalvanized steel according to claim 30, wherein the steel sheet isreplaced with a steel sheet having a composition comprising, in weightpercent, C: not more than 0.20%, Si: not more than 2.0%, Mn: not morethan 3.0%, P: not more than 0.1%, S: not more than 0.02%, Al: not morethan 0.3%, N: not more than 0.02%, and at least one selected from thegroup consisting of Mo: 0.05 to 2.0%, Cr: 0.05 to 2.0% and W: 0.05 to2.0% in a total amount of not more than 2.0%.
 33. A method formanufacturing a high-ductility hot-dip galvanized steel sheet accordingto claim 32, the composition further containing, in weight percent, atleast one of Nb, Ti, and V in a total amount of not more than 2.0%. 34.A method for manufacturing a high-ductility hot-dip galvanized steelsheet according to claim 30, further comprising a pickling treatmentstep of pickling the steel sheet between the primary heat treatment stepand the secondary heat treatment step.
 35. A method for manufacturing ahigh-ductility hot-dip galvanized steel sheet according to claim 30,further comprising an alloying step of alloying the hot-dip galvanizinglayer, subsequent to the hot-dip galvanizing step.
 36. A method formanufacturing a high-strength hot-dip galvanized steel sheet accordingto claim 30, wherein the steel sheet is a hot rolled steel sheetmanufactured by hot-rolling a material under conditions including aheating temperature of not less than 900° C., a finish rolling endtemperature of not less than 700° C. and a coiling temperature of notmore than 800° C., or a cold-rolled steel sheet obtained by cold-rollingthe hot-rolled steel sheet.
 37. A method for manufacturing ahigh-strength hot-dip galvanized steel sheet according to claim 36,wherein the cool-rolling is performed at a reduction ratio of not lessthan 40%.